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Handbook of Conducting Polymers Third Edition
CONJUGATED POLYMERS PROCESSING AND APPLICATIONS
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Handbook of Conducting Polymers Third Edition
CONJUGATED POLYMERS PROCESSING AND APPLICATIONS
Edited by
Terje A. Skotheim and John R. Reynolds
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CRC Press Taylor & Francis Group 6000 Broken Sound Parkway NW, Suite 300 Boca Raton, FL 33487-2742 © 2007 by Taylor & Francis Group, LLC CRC Press is an imprint of Taylor & Francis Group, an Informa business No claim to original U.S. Government works Printed in the United States of America on acid-free paper 10 9 8 7 6 5 4 3 2 1 International Standard Book Number-10: 1-4200-4360-9 (Hardcover) International Standard Book Number-13: 978-1-4200-4360-0 (Hardcover) This book contains information obtained from authentic and highly regarded sources. Reprinted material is quoted with permission, and sources are indicated. A wide variety of references are listed. Reasonable efforts have been made to publish reliable data and information, but the author and the publisher cannot assume responsibility for the validity of all materials or for the consequences of their use. No part of this book may be reprinted, reproduced, transmitted, or utilized in any form by any electronic, mechanical, or other means, now known or hereafter invented, including photocopying, microfilming, and recording, or in any information storage or retrieval system, without written permission from the publishers. For permission to photocopy or use material electronically from this work, please access www.copyright.com (http:// www.copyright.com/) or contact the Copyright Clearance Center, Inc. (CCC) 222 Rosewood Drive, Danvers, MA 01923, 978-750-8400. CCC is a not-for-profit organization that provides licenses and registration for a variety of users. For organizations that have been granted a photocopy license by the CCC, a separate system of payment has been arranged. Trademark Notice: Product or corporate names may be trademarks or registered trademarks, and are used only for identification and explanation without intent to infringe. Library of Congress Cataloging-in-Publication Data Skotheim, Terje A., 1949Conjugated polymers: processing and applications. -- 3rd ed. / Terje A. Skotheim, John Reynolds. p. cm. Includes bibliographical references and index. Previously published: Handbook of conducting polymers. 2nd ed., rev. and expanded. c1998. ISBN-13: 978-1-4200-4360-0 ISBN-10: 1-4200-4360-9 1. Conducting polymers. 2. Organic conductors. I. Reynolds, John R., 1956- II. Skotheim, Terje A., 1949- Handbook of conducting polymers. III. Title. QD382.C66H36 2006b 620.1’92--dc22 Visit the Taylor & Francis Web site at http://www.taylorandfrancis.com and the CRC Press Web site at http://www.crcpress.com
2006036331
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Dedication
This book is dedicated to our spouses, Ellen Skotheim and Dianne Reynolds. Without their understanding and support, we would never have completed this project.
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Preface to Third Edition
The field of conjugated, electrically conducting, and electroactive polymers continues to grow. Since the publication of the second edition of the Handbook of Conducting Polymers in 1998, we have witnessed broad advances with significant developments in both fundamental understanding and applications, some of which are already reaching the marketplace. It was particularly rewarding to see that in 2000, the Nobel Prize in chemistry was awarded to Alan Heeger, Alan MacDiarmid, and Hideki Shirakawa, recognizing their pathbreaking discovery of high conductivity in polyacetylene in 1977. This capstone to the field was celebrated by all of us as the entire community has participated in turning their initial discovery into the important field that it now is, almost 30 years later. The vast portfolio of new polymer structures with unique and tailored properties and the wide range of applications being pursued are far beyond what we could have envisioned when the field was in its infancy. It was developments in polymer synthesis that led to free-standing polyacetylene films and the discovery of conductivity in polymers. The synthesis of p-conjugated chains is central to the science and technology of conducting polymers and is featured in this edition. Examining the synthetic advances across the board, one is struck by refined and careful syntheses that have yielded polymers with well-controlled and wellunderstood structures. Among other things, it has led to materials that are highly processable using industrially relevant techniques. In aspects of processing spin coating, layer-by-layer assembly, fiber spinning, and the application of printing technology have all had a big impact during the last 10 years. Throughout the Handbook, we notice that structure–property relationships are now understood and have been developed for many of the polymers. These properties span the redox, interfacial, electrical, and optical phenomena that are unique to this class of materials. During the last 10 years, we have witnessed fascinating developments of a wide range of commercial applications, in particular, in optoelectronic devices. Importantly, a number of polymers and compositions have been made available by the producers for product development. This has helped to drive the applications developments to marketable products. While conductivity, nonlinear optics, and light emission continue to be important properties for investigation and have undergone significant developments as discussed throughout the Handbook, the advances in semiconducting electronics, memory materials, photovoltaics (solar cells), and applications directed to biomedicine are emerging as future growth areas. As we have assembled this edition, it has become clear that the field has reached a new level of maturity. Nevertheless, with the vast repertoire of synthetic chemistry at our disposal to create new structures with new, and perhaps unpredictable properties, we can expect exciting discoveries to continue in this dynamic field.
Terje A. Skotheim Tucson, Arizona John R. Reynolds Gainesville, Florida
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Editors
Terje Skotheim is the founder and Chief Executive of Intex, a Tucson, Arizona technology company. Dr. Skotheim is an experienced developer of several technologies, a seasoned executive and a successful founder of several startup companies in the United States, Norway, and Russia. His research interests over more than 25 years span several disciplines in materials science and applications, including electroactive and conjugated polymers, molecular electronic materials, solid-state ion conductors, new electronic nanoamorphous carbon- and diamond-like carbon materials, and thin-film and surface science. He has pursued a wide range of technology applications of advanced materials in OLEDs, biosensors, lithium batteries, photovoltaic cells, and MEMS devices. He has held research positions in France, Sweden, and Norway in addition to the United States and was head of the conducting polymer group at Brookhaven National Laboratory before launching his career as an entrepreneur. Skotheim received his B.S. in Physics from the Massachusetts Institute Technology and Ph.D. in physics from the University of California at Berkeley (1979). He is the editor/co-editor of the Handbook of Conducting Polymers (1st and 2nd editions, Marcel Dekker) and Electroresponsive Molecular and Polymeric Systems (Marcel Dekker), the author of more than 300 publications and more than 70 patents. He can be reached at
[email protected] John R. Reynolds is a professor of chemistry at the University of Florida with expertise in polymer chemistry. He serves as an associate director for the Center for Macromolecular Science and Engineering. His research interests have involved electrically conducting and electroactive conjugated polymers for over 25 years, with work focused on the development of new polymers by manipulating their fundamental organic structure in order to control their optoelectronic and redox properties. His group has been heavily involved in the areas developing new polyheterocycles, visible and infrared light electrochromism, along with light emission from polymer and composite light-emitting diodes (LEDs) (both visible and near-infrared) and light emitting electrochemical cells (LECs). Further work is directed to using organic polymers and oligomers in photovoltaic cells. Reynolds obtained his M.S. (1982) and Ph.D. (1984) in polymer science and engineering from the University of Massachusetts. He has published over 200 peer-reviewed scientific papers and served as co-editor of the Handbook of Conducting Polymers, which was published in 1998. He can be reached by e-mail at
[email protected] or visit his Web site http://www.chem.ufl.edu/~reynolds/.
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Contributors
Michal Adler Department of Chemistry Solid State Institute Technion Haifa, Israel
Mary Galvin Department of Materials Science and Engineering University of Delaware Newark, Delaware
Magnus Berggren Dept. of Science and Technology Linko¨ping University Norrko¨ping, Sweden
Elena Gershman Department of Chemistry Solid State Institute Technion Haifa, Israel
Gordon P. Bierwagen Department of Coatings and Polymeric Materials North Dakota State University Fargo, North Dakota Sean Brahim Center for Bioelectronics, Biosensors, and Biochips Clemson University Clemson, South Carolina Hermona Christian-Pandya Department of Biomedical Engineering Northwestern University Evanston, Illinois Chih-Wei Chu Department of Materials Science and Engineering University of California Los Angeles, California Larry R. Dalton Department of Chemistry University of Washington Washington, D.C. Yoav Eichen Department of Chemistry Solid State Institute Technion Haifa, Israel
Oded Globerman Nanoelectronic Center Electrical Engineering Department Technion Haifa, Israel Vladimir Gorelik Department of Chemistry Solid State Institute Technion Haifa, Israel Anthony Guiseppi-Elie Center for Bioelectronics, Biosensors, and Biochips (C3B) and Department of Chemical and Biomolecular Engineering and Department of Bioengineering Clemson University Clemson, South Carolina and Abtech Scientific, Inc. Biotechnology Research Park Richmond, Virginia Olle Ingana¨s Biomolecular and Organic Electronics, IFM Linko¨ping University Linko¨ping, Sweden
Terje A. Skotheim / Conjugated Polymers: Processing and Applications
P.C. Innis ARC Centre of Excellence for Electromaterials Science Intelligent Polymer Research Institute University of Wollongong Wollongong, Australia David J. Irvin Naval Air Warfare Center Weapons Division Chemistry Division China Lake, California Jennifer A. Irvin Naval Air Warfare Center Weapons Division Chemistry Division China Lake, California Ghassan E. Jabbour Flexible Display Center and Department of Chemical and Materials Engineering Arizona State University Tempe, Arizona Benjamin R. Mattes Santa Fe Science and Technology University of New Mexico Santa Fe, New Mexico S.E. Moulton ARC Centre of Excellence for Electromaterials Science Intelligent Polymer Research Institute University of Wollongong Wollongong, Australia A.J. Mozer Molecular Process Engineering, Material and Life Science Graduate School of Engineering Osaka, Japan and
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Ian D. Norris Santa Fe Science and Technology University of New Mexico Santa Fe, New Mexico Toribio F. Otero Center for Electrochemistry and Intelligent Materials Universidad Polite´cnica de Cartagena Cartagena, Spain Jianyong Ouyang Department of Materials Science and Engineering University of California Los Angeles, California Ankita Prakash Department of Materials Science and Engineering University of California Los Angeles, California Yevgeni Preezant Nanoelectronic Center Electrical Engineering Department Technion Haifa, Israel Noam Rappaport Nanoelectronic Center Electrical Engineering Department Technion Haifa, Israel Nathaniel D. Robinson Department of Science and Technology Linko¨ping University Norrko¨ping, Sweden Yohai Roichman Nanoelectronic Center Electrical Engineering Department Technion Haifa, Israel
Linz Institute for Organic Solar Cells Physical Chemistry Johannes Kepler University Linz Linz, Austria
N.S. Sariciftci Linz Institute for Organic Solar Cells Physical Chemistry Johannes Kepler University Linz Linz, Austria
Peter Nilsson Biomolecular and Organic Electronics, IFM Linko¨ping University Linko¨ping, Sweden
Elisabeth Smela Department of Mechanical Engineering University of Maryland College Park, Maryland
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Olga Solomesch Nanoelectronic Center Electrical Engineering Department Technion Haifa, Israel Geoffrey M. Spinks ARC Centre for Electromaterials Science Intelligent Polymer Research Institute University of Wollongong Wollongong, Australia John D. Stenger-Smith Naval Air Warfare Center Weapons Division Chemistry Division China Lake, California Shay Tal Department of Chemistry Solid State Institute Technion Haifa, Israel Dennis E. Tallman Department of Chemistry, Biochemistry and Molecular Biology and Department of Coatings and Polymeric Materials North Dakota State University Fargo, North Dakota Nir Tessler Nanoelectronic Center Electrical Engineering Department Technion Haifa, Israel Ricky J. Tseng Department of Materials Science and Engineering University of California Los Angeles, California Subramanian Vaidyanathan Solutions Polymers Research—Energy and Electronics BASF AG Ludwigshafen, Germany
Janos Veres Merck Chemicals Manchester, United Kingdom G.G. Wallace ARC Centre of Excellence for Electromaterials Science Intelligent Polymer Research Institute University of Wollongong Wollongong, Australia Bernard Wessling Ormecon International Ormecon GmbH Ammersbek, Germany Ann M. Wilson Abtech Scientific, Inc. Biotechnology Research Park Richmond, Virginia Yang Yang Department of Materials Science and Engineering University of California Los Angeles, California Yuka Yoshioka Flexible Display Center and Department of Chemical and Materials Engineering Arizona State University Tempe, Arizona Vadim Zolotarev Department of Chemistry Solid State Institute Technion Haifa, Israel
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Table of Contents
I:
Processing of Conjugated Polymers
1.
Conductive Polymers as Organic Nanometals ........................................................................... 1-3 Bernard Wessling
2.
Conducting Polymer Fiber Production and Applications......................................................... 2-1 Ian D. Norris and Benjamin R. Mattes
3.
Inkjet Printing and Patterning of PEDOT–PSS: Application to Optoelectronic Devices............................................................................................................ 3-1 Yuka Yoshioka and Ghassan E. Jabbour
4.
Printing Organic Electronics on Flexible Substrates................................................................. 4-1 Nathaniel D. Robinson and Magnus Berggren
II:
Applications and Devices Based on Conjugated Polymers
5.
Polymers for Use in Polymeric Light-Emitting Diodes: Structure–Property Relationships ............................................................................................... 5-3 Hermona Christian-Pandya, Subramanian Vaidyanathan, and Mary Galvin
6.
Organic Electro-Optic Materials ................................................................................................. 6-1 Larry R. Dalton
7.
Conjugated Polymer Electronics—Engineering Materials and Devices................................... 7-1 Nir Tessler, Janos Veres, Oded Globerman, Noam Rappaport, Yevgeni Preezant, Yohai Roichman, Olga Solomesch, Shay Tal, Elena Gershman, Michal Adler, Vadim Zolotarev, Vladimir Gorelik, and Yoav Eichen
8.
Electrical Bistable Polymer Films and Their Applications in Memory Devices ..................... 8-1 Jianyong Ouyang, Chih-Wei Chu, Ricky J. Tseng, Ankita Prakash, and Yang Yang
9.
Electroactive Polymers for Batteries and Supercapacitors........................................................ 9-1 Jennifer A. Irvin, David J. Irvin, and John D. Stenger-Smith
10.
Conjugated Polymer-Based Photovoltaic Devices.................................................................... 10-1 A.J. Mozer and N.S. Sariciftci
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11.
Biomedical Applications of Inherently Conducting Polymers (ICPs).................................... 11-1 P.C. Innis, S.E. Moulton, and G.G. Wallace
12.
Biosensors Based on Conducting Electroactive Polymers....................................................... 12-1 Anthony Guiseppi-Elie, Sean Brahim, and Ann M. Wilson
13.
Optical Biosensors Based on Conjugated Polymers ................................................................ 13-1 Peter Nilsson and Olle Ingana¨s
14.
Conjugated Polymers for Microelectromechanical and Other Microdevices ........................ 14-1 Geoffrey M. Spinks and Elisabeth Smela
15.
Corrosion Protection Using Conducting Polymers ................................................................. 15-1 Dennis E. Tallman and Gordon P. Bierwagen
16.
Artificial Muscles ........................................................................................................................ 16-1 Toribio F. Otero
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I Processing of Conjugated Polymers
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1 Conductive Polymers as Organic Nanometals 1.1
Organic Metals..................................................................... 1-3 Metallic Character and Nanostructure of Conductive Polymers . Nanotechnology with Nanometals
1.2 1.3
Bernard Wessling
1.1
Conductive Polymers–Solvent Systems: Solutions or Dispersions................................................... 1-22 Applications of Organic Metals Emerging from Basic Science: Macroeffects of Nanoparticles ................. 1-23 Introduction Applications
. .
PAni Dispersions and Blends Final Remarks
Organic Metals
The broad potential and the scientific and technical possibilities of organic metals are widely unknown. An interesting publication in Nature shows how scientists are oriented toward nanotechnology with conventional metals, even if the approach is more complicated. In their research work, Erez Braun et al. [1] placed a DNA double strand between the two electrodes to be contacted, and then deposited about 30–50 nm silver particles by the reduction of a Agþ solution on the DNA. When only 10–20 nm silver particles were deposited, they did not make successful electrical contact. This is not surprising in view of the fact that conventional nanometals, like Ag or Au, do not exhibit a very high conductivity, as evident by their conducitivity values of only <104–106 S=cm. Only annealing at rather high temperature (>1008C), which destroys the nano character, increases the conductivity to the metallic regime. In contrast, the existence of organic metals and their character as true metals, though nanometals, is still widely unknown although they show a much higher conductivity, in spite of their nanosize, after deposition from colloidal dispersion (between 1 and <102 S=cm). ‘‘Conductive polymers,’’ in contrast, is a well-known term [2]. Polyacetylene, polyaniline (PAni), polypyrrole, polythiophene, and many more polymers have been synthesized and studied. Under the headline ‘‘conductive polymers,’’ however, the possibilities of studying nanoscale phenomena are not broadly considered, as those groups in the scientific community focusing on their polymeric character are trying to treat them primarily like other polymers (though conductive) and related their properties to the individual chains (Figure 1.1) instead of looking at them primarily as potential metals, organic metals, or synthetic nanometals. The difference in the understanding is centered around the question of whether the primary structural and functional unit for the conductive polymers is the single (eventually oriented) chain, and the main 1-3
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Conjugated Polymers: Processing and Applications
Rod-like
Expanded coil-like
Coil-like
Metallic fibril
1d Weak link
(a)
(b)
FIGURE 1.1 (a) Present understanding of conductivity between metallic islands (or metallic fibrils, connected by amorphous individual chains. (Reprinted from MacDiarmid, A.G. and Epstein A.J., Synth. Met., 65, 103, 1994. With permission. Copyright 1994 Elsevier Science). (b) Comparison of two different interpretations of polyacetylene fibrils: flat platelets spherical globules. (Reprinted from Wessling, B., Makromol. Chem., 185, 1265, 1984. With permission. Copyright 1984 Wiley–VCH Verlag GmbH)
transport mechanism is the transport of polarons, or the primary unit of the organic metal is a nanoparticle of 10 nm and the transport mechanism is metallic plus tunneling (from particle to particle). Here again, it is very interesting to see that even in reviews under the heading of ‘‘Metal Clusters and Colloids’’ [3], organic metals and their colloidal and nanostructural properties are not mentioned. We will therefore discuss the nanoparticle structure and dynamics, the generation of the metallic character, the precondition for these properties and for eventual nanotechnology, the principal insolubility of organic metals (and of conductive polymers as well), and some actual technical applications ultimately based on their nanocharacter and on dispersion and the macrotechnology with dramatic effects in the nanoscale.
1.1.1
Metallic Character and Nanostructure of Conductive Polymers
The deciding hint for explaining the basic electron transport mechanism in conductive polymers came from studies with nanoparticles of conventional metals like indium, silver, or copper. Nimtz et al. [4] found in 1989 that metallic particles, if prepared on a nanoscale of below 1 mm down to 10 nm, show some distinct deviations from macroscopic metals (see Figure 1.2). For the first time, Nimtz succeeded to prove experimentally that the conductivity of regular metals, if they are prepared in mesoscopic form (i.e., in a size between macro- and microscopic: hence nanometals), differs basically from that of macroscopic metals. Not only is the conductivity dependance on temperature no longer purely metallic, but also decreases with the decreasing temperature.
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1-5
Conductive Polymers as Organic Nanometals
108 Bulk
106 6 (Ωm)−1 104
Classical d3 Quantum Experiment
102
100
Indium (T = 300 k)
10−2 10−8
10−7
10−6
10−5
d (m)
FIGURE 1.2 Principal behavior of size-dependent conductivity according to classical or quantum theoretical interpretation, compared with experiment. (Reprinted from Nimtz, G., Enders, A., Marquardt, P., Pelster, R., and Wessling, B., Synth. Met., 45, 197, 1991. With permission. Copyright 1991 Elsevier Science)
Normally, the metallic conduction band extends over macroscopic distances and allows electrons (the electron gas) to move freely, only interacting with phonons (lattice vibrations). With decreasing temperature, the vibrations lose intensity, and hence the electrons can move even better, which is the basis for the increasing conductivity of metals with decreasing temperature. In nanometals, however, the conduction band has a size in the same range as the electron wavelengths. Therefore, only certain wavelengths are allowed, i.e., those that have a node plane at the boundaries of the three-dimensional conduction band. Hence, the conductivity is quantum size limited. We observe the quantum effects of conductivity. These results motivated us to work together with the Cologne group and to find out whether there are similarities between mesoscopic metals and organic metals [5], as we suspected that the conductivity phenomena in conductive polymers may be better understood taking nanostructures into account. 1.1.1.1 Nanoparticulate Morphology and Dispersion The background of the idea was our findings in previous work that conductive polymers are composed of more or less globular primary particles (Figure 1.3). We first found [6] that polyacetylene, even though apparently in a fibrillar morphology (as concluded from transmission electron microscopy), clearly showed a particle substructure with a particle size of 100 nm (Figure 1.4). Even at this early time, we saw fine structures below 100 nm, but we had not yet been able to conclude that the particle size was even finer. Even at a 100 nm size, one would have to consider conductive polymers as being made up of nanoparticles. We will see further in the chapter that these (secondary) particles play an important role in the dispersion process in polymeric media. We constructed a model (Figure 1.1a) according to which the fibrillar morphology, which was seen by many researchers, is due to an oriented arrangement of these particles in a pearl-chain-like structure. We
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23892 (d) 23892B 25385
(a–c)
(e)
FIGURE 1.3 Scanning electron microscopic evidence for globular or spherical subunits in fibrils detected by TEM. (Reprinted from Wessling, B., Makromol. Chem., 185, 1265, 1984. With permission. Copyright 1984 Wiley–VCH Verlag GmbH)
FIGURE 1.4 TEM of flocculated polyaniline in polymer matrix, showing further fine structure in the about 100 nm big secondary particles. (Reprinted from (a) Genie`s, E., Intrinsically Conductive Polymers–An Emerging Technology, ed. M. Aldissi, Kluwer Academic, Dordrecht, the Netherlands, 1993, 75; (b) Genie`s, E. New J. Chem., 15, 373, 1991. With permission. Copyright 1997 John Wiley & Sons Limited)
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1-7
Conductive Polymers as Organic Nanometals
% Pass 100.0
% Chan 20.0
90.0
18.0
80.0
16.0
70.0
14.0
60.0
12.0
50.0
10.0
40.0
8.0
30.0
6.0
20.0
4.0
10.0
2.0
0.0 0.0010
0.0100
0.1000 - Size (µm) -
0.0 10.00
1.000
FIGURE 1.5 Laser Doppler measurement of polyaniline primary particles in an organic solvent dispersion. (Reprinted from Wessling, B., Handbook of Nanostructured Materials and Nanotechnology, vol. 5, ed. H.S. Nalwa, Academic Press, New York, 1999, 525. With permission. copyright 2000 Elsevier Science)
later found [7] by membrane filtration that the basic primary particle in obviously all conductive polymers is 10 nm in size. Subsequently, in more intensive studies, using various techniques like scanning tunneling microscopy (STM) [8] and photon correlation spectroscopy [9], and routinely via laser Doppler measurements [10] (Figure 1.5), we confirmed a primary particle size of 10 nm. Independent of our research, STM studies [11] on the morphology of oriented polyacetylene with very high conductivity (so-called N-PAc [12]) revealed very densely aggregated particular subunits in the fibrillar structure with maximum contact area between them (Figure 1.6). With the first insight into the particulate morphological substructure, our very early concept of processing conductive polymers via dispersion was supported for the first time. We decided to polymerize powders, with preferably very well-displayed globular morphology, and we were successful with our first dispersion (Figure 1.7). At that time, it was rather unclear how and at which concentration a dispersed conductive polymer would conduct after being dispersed in a (polymeric) matrix. With our first patent application [13], we
1 µm
a e
f
g
0.1 µm
0.1 µm
00.5 µm
0.5 µm
b
0.25 µm 0.1 µm
c
d
FIGURE 1.6 Primary particles in stretch aligned, highly conducting polyacetylene. (Reprinted from Theophilou, N., Swanson, D.B., MacDiarmid, A.G., Mantovani, J.G., Annis, B.K., and Epstein, A.J, Electronic Properties of Conjugated Polymers III (Springer series in solid-state sciences, vol. 91), eds. H. Kuzmany, M. Mehring, and S. Roth, Springer– Verlag, Berlin, 1989, 14–18. With permission. Copyright 1989 Springer–Verlag)
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Conjugated Polymers: Processing and Applications
FIGURE 1.7 Secondary particles of polyacetylene (a) the powder as synthesized before being dispersed in a polymer matrix at 0.1% (b). (Part (a) is reprinted from Wessling, B., Handbook of Nanostructured Materials and Nanotechnology, vol. 5, ed. H.S. Nalwa, Academic Press, New York, 1999, 525. With permission. Copyright 2000 Elsevier Science; Part (b) is reprinted from Wessling, B. and Volk, H, Synth. Met., 15, 183, 1986. With permission. Copyright 1986 Elsevier Science.)
had published that doped PAc, after dispersion, caused a conductivity breakthrough for the whole system at a critical volume concentration around 10%. On the basis of the picture of fibrillar morphology in which chains oriented in the fibril direction are the structural basis for solitonic or polaronic electron transport mechanisms, such a result would only be understandable if the fibrils were arranged in the matrix in a stretched, unoriented, nonparallel form, which was definitely not the case. Below the critical volume concentration, we found particles (instead of fibrils) isolated from each other, in a size and a form that can also be detected as the (secondary) morphological units of the raw powder (see Figure 1.7). But this does not explain how the conductivity would rise at a certain critical concentration or why this happens (with a drastic jump of many orders of magnitude) at a certain concentration. Carbon black or metal powder containing polymer compounds show a similar behavior when dispersed in a matrix above certain different critical concentrations. The percolation theory is thought to be the best tool for the description of this effect [14]. It is believed that metal powder, having a globular particle shape, is distributed in a statistically even manner and the powder particles will make contacts, governed by statistical laws (probability), whenever enough particles are present and close enough to finally form the first continuous conductive pathways. The critical volume concentration of metal powders is within the range of what is predicted by the percolation theory (45–64 vol%). For carbon black, however, the theory cannot explain the partially rather low critical concentrations, between 25 and 10 vol% and in well-defined optimal cases even down to 1% and below. Percolation theorists assume that carbon black particles are highly structured, with a high length-to-diameter ratio of their arms, and therefore have a bigger chance of contacting each other. At and above the critical concentration, a sudden change in the arrangement of the particles occurs: the previously well-dispersed and well-separated particles form complex networks. We found that dispersion led to a rather complex arrangement of phases, adsorbed layers, and finally even more complex flocculation structures in form of networks. Within the networks (Figure 1.8), the particles can touch and at least contact the next neighbors. The three-dimensional connectivity of the twodimensional networks is being provided by the further complex three-dimensional arrangements and structures of the dispersion and flocculation layers. Dispersion, although is a process pushed by macroscopic tools and processes, results in very fine and precisely exhibited nanostructures: in a polymeric matrix, a monomolecular layer of 15 nm thickness is forced to adsorb on the particles becoming dispersed down to a particle size of 50–250 nm (depending on matrix nature, process efficiency etc.). These particles (with their adsorbed layer) phase separate to monolayers (with a thickness of 80–280 nm), wherein flocculation and network formation occurs at and
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1-9
Conductive Polymers as Organic Nanometals
1
2
A
A
3.1
3.2
4
5
B
B
00604 300AM
00628 300AM (a)
(b)
(c)
Matrix
Matrix Seams A
Seams Seams and matrix both More realistic form: seams curved continuous phases interpenetrating network, simplified representation
(d)
Matrix Seams B
Seams
(e)
FIGURE 1.8 SEM studies on dispersed and flocculated conductive phases and mechanical interpretation model for the flocculation process. (Part (a) is reprinted from Wessling, B., Polym. Eng. Sci., 31, 1200, 1991. With permission. Copyright 1991 Polymer Science and Engineering; Part (b) is reprinted from Wessling, B., Synth. Met., 45, 119, 1991. With permission. Copyright 1991 Elsevier Science.)
above the critical volume concentration, with branching roughly every 10 nanoparticles, building a network size comprising of 30–50 particles (in one network element). Such structures, although chaotic and based on nonequilibrium thermodynamics [15], are nonetheless reproducibly formed. Reproducibly meaning that the nanoscopic properties leading to macroscopic properties can be reproducibly achieved, although the responsible nanostructures are widely differing and never identical. Their formation is purely a process of self-organization [16]. And although the resulting exact structures are never the same, it is the inherent nature of the self-organization process on the nanoscale that allows stability and reproducibility. In an analogous way, dispersion in nonpolymeric media (solvents) can be understood. The first difference is that we can now disperse down to the primary particle size, 10 nm (cf. particle size analysis with laser Doppler, Figure 1.5), because no solid adsorbed layer is formed. In contrast, an adsorbed layer of solvent molecules is formed on the surface of the particles being dispersed. Interestingly, at very low concentrations such dispersions suddenly gel, probably because the dispersed particles arrange into very fine pearl-chains like filaments, or even more complex threedimensional branched tubes (Figure 1.9c), which is the reason for the photon correlation peak around 250 nm (Figure 1.9b). Again, this is a self-organization process of nanoparticles under nonequilibrium conditions (see Figure 1.10).
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100
100
90
90
80
80
70
70
60
60
Intensity
Intensity
Conjugated Polymers: Processing and Applications
50 40
30
20
20
10
10 0
10 100 1 Particle diameter (nm)
1000 10000 100 Particle diameter (nm)
100
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Intensity
Intensity
40
30
0
50 40
50 40
30
30
20
20
10
10
0 (a)
50
0 10 100 1 Particle diameter (nm)
1000 10000 100 Particle diameter (nm)
(b)
FIGURE 1.9 Photon correlation spectroscopy of primary particle dispersions of neutral and doped polyaniline; cf Figure 1.5, where laser Doppler spectroscopy only shows the true particle diameter, whereby photon correlation spectroscopy also shows superstructures (Part (a) is reprinted from Wessling, B., Handbook of Nanostructured Materials and Nanotechnology, vol. 5, ed. H.S. Nalwa, Academic Press, New York, 1999, 525. With permission. Copyright 2001 Elsevier Science; Part (b) is reprinted from Wessling, B., Adv. Mater., 5, 300, 1993. With permission. Copyright Wiley–VCH Verlag GmbH.)
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Conductive Polymers as Organic Nanometals
FIGURE 1.10 SEM freeze drying. (Reprinted from Wessling, B., Handbook of Nanostructured Materials and Nanotechnology, vol. 5, ed. H.S. Nalwa, Academic Press, New York, 1999, 525. With permission. Copyright 2000 Elsevier Science)
1.1.1.2 Metallic Properties and Conduction Mechanism of the Organic Nanometal Now that we know of the particulate morphology and the complex dynamics under which the particles become dispersed and self-organize to continuous complex networks and filaments, it remains to be discussed how conductivity is achieved in a dispersion (above the critical concentration) (Figure 1.11).
1E+2 PAni PMMA #900210 (1995)
1E+0
PAni #172 (1988)
Conductivity S/cm
1E−2
PAni (1986)
1E−4 1E−6
PAc (1984)
1E−8 1E−10 1E−12 0
10
20
30
40
50
Concentration vol.%
FIGURE 1.11
Critical concentration for various generations of organic metal dispertibility between 1986 and 1995.
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We will start with the observations: At the critical concentration, conductivity jumps over several orders of magnitude, and the flocculated network structures can be observed under the scanning electron microscope. Many theories have been developed (involving solitons, excitons, polarons and bipolarons) [17] to explain the conductivity phenomenon under the assumption that the chains of conductive polymers are being arranged and at least somewhat oriented in fibrils. But now, it must be explained why our dispersed (and later flocculated) polymer showed principally the same transport properties as the fibrillar conductive polymers, as can be concluded from conductivity versus temperature and thermopower measurements. The thermopower measurements themselves show that dispersed conductive polymers have essentially the same thermopower behavior as metals: it is small (semiconductors show a high S), and decreases practically linearly with decreasing temperature [18]. These properties were known from our previous measurements [19], which showed the linear temperature dependance for a range of 100 K. These results lead us to the assumption that the primary transport mechanism is metallic, but obviously influenced by a barrier mechanism. To overcome the barrier, a temperature-activated transport process had to be active as well. Our cooperation with Guenther Nimtz of Cologne University showed very quickly [5] that the conductive polymer (in the first experiments deposited from dispersions onto PC films) behaves principally in the same way as Nimtz’s mesoscopic metals, which he had prepared by vapor deposition in oil to form a colloidal dispersion of finest metal droplets in oil. But whereas conventional metals are hard to prepare in a small particle size, and their quantum size effect is hard to observe, the conductive polymers only occur (at least according to our results and conclusions)—due to their extremely high surface tension—in nanoparticulate morphology. So we should see their quantum size-limited conductivity rather easily, and this is indeed the case. Rolf Pelster studied the quantum effects in close cooperation with us in various dispersions and in the pure polyaniline [20]. We found that the primary metallic unit is 8 nm in size, surrounded by a nonmetallic (amorphous) layer of the same composition, but with a less optimal structure (so that a particle size of 9.6 nm was determined, which is very close to the 10 nm value that we had earlier found by other techniques). This and the dielectric between the contacting particles are the barriers through which the electron wave (thermally activated) will eventually tunnel to reach the neighboring particles (Figure 1.12). These results forced us to conclude that the conductive polymers are in fact organic metals or they are nanometals, and two different transport mechanisms contribute to the conduction mechanism: (a) a purely metallic part within each particle and (b) a thermally activated part from one particle to another (Figure 1.12b). This explains why there is no principal difference between the raw (doped) polyaniline and its dispersions, in whatever medium. The differences to be observed were only of quantitative, and not of qualitative nature, at least not in the direction, which was expected by most of those who still favor the fibril hypothesis. They believe that the chain is the primary active unit, which could also be dissolved and is believed to have conductive properties, even as a single chain. If that were the case, a dispersed (i.e., mechanically separated, in case of assumed fibrillar morphology, even destroyed) conductive polymer would not have the same conductivity and especially not the same transport properties in the dispersion medium above the critical volume concentration or after deposition from a low molecular weight liquid medium and drying. Such results (i.e., lower conductivity in blends and at maximum concentrations, compared to the pure conductive polymer) can often be found in the literature [21]. Our results showed the contrary (Figure 1.13): we observed the lowest conductivity at room temperature and at concentrations between 25% and 40% (i.e., 18–30 vol% only). This conductivity was in the range of the conductivity of the undispersed raw material: 5–10 S=cm [22]. In the best dispersions at 40%, we find around 50–100 S=cm, 10 times more than the undispersed pure raw material. At any temperature lower than
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ζ
(a)
(b)
FIGURE 1.12 Model of electron wave density distribution, barrier heights, and crystalline core in the primary particles, and visualization of the limitation of electron wavelength in the nanometallic primary particles to such values that have a node plane at the particle walls. (Reprinted from Wessling, B., Handbook of Organic Conductive Molecules and Polymers, vol. 3, ed. H.S. Nalwa, Wiley, Chichester, 1997, 497–632. With permission. Copyright 1997 John Wiley & Sons Limited)
the ambient temperature, the blends were several orders of magnitude more conductive than the pressed powder. This shows that the metallic contribution to the conductivity is the same in the raw undispersed organic metal; the temperature activated tunneling process, however, is easier in the dispersion—or, better, after dispersion!
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1.2 PAni blends 1
60% PMMA
s (T ) / s (300)
0.8
53% PVC 100% PAni
0.6
60% PETG
0.4
0.2
0 0
100
(a)
T (K)
200
300
10 67% PMMA 60% PETG 60% PMMA
8
6 S (µV K−1)
PETG blends 4 PAni 2
0
−2 0
100
200
300
T (K) (b)
FIGURE 1.13 Transformation of a pure premetallic organic metal (100% PAni) to a truly metallic polyaniline in blends after dispersions. (Reprinted from Gospodinova, N., Mokreva, P., Tsanov, T., and Terlemeyzan, L., Polymer, 38, 743, 1997. With permission. Copyright 1996 Elsevier Science)
In contrast to expectations of most scientists, dispersion does not lead to a deterioration of the conductivity properties (by, e.g., new barriers from the dispersion matrix), but to an improvement. This implies that during the dispersion process, barriers must have been reduced, and possibly the arrangement of the particles in relation to each other (in the flocculated network) is optimized with regard to the tunneling process.
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So, it was only a small surprise when we found a 10 times higher conductivity (between 20 and 100 S=cm) at 40% organic metal content in a thermoplastic polyacrylic dispersion. Furthermore, for the first 70 K temperature decrease, conductivity increased, as is the case for conventional metals (Figure 1.13) [23]. A new effect had been discovered: the raw organic metal powder, undispersed and pressed to a pellet, has the normal conductivity of 5 S=cm and a nonmetallic temperature dependance of conductivity (but a metallic thermopower). When dispersed (and even at lower concentrations such as 30 vol%), it has an absolutely higher conductivity at ambient temperature, a partially metallic conductivity or temperature dependance and the conductivity at 10 K decreased only to 5–10 S=cm, whereas the starting raw material has a conductivity of 1010 S=cm (cf Figure. 1.13a) [23]. An earlier assumption (at that time explaining the fibrillar secondary morphology of conductive polymers), according to which counterions are anisotropically arranged in the primary particles that leads to a principally linear arrangement of the particles (Figure 1.14), also could explain why the additional dispersion process would result in an increase of the conductivity: a preferred linear arrangement of the particles may have been caused. Further proof for the (in the worst case) unchanged or even improved metallic character of the organic metal after dispersion came from ESR and magnetoconductivity measurements [24]. First studies, performed on polyaniline–polyester blends (the same as used in Ref. [24]), showed a much higher Pauli susceptibility and much higher density of states at the Fermi level as did even the relatively highly conducting polyaniline–camphersulfonic acid as doped in m-cresol [25]. Until then, it was a common opinion that only solution processing (also often called secondary doping [26]) could improve the conductivity, and these effects were related to better orientation of the chains. The scientific community did not yet recognize the importance of these results for the understanding of the conductive polymers as organic metals. Further studies [24] showed that according to the reduced activation energy, W versus T log–log plot between 300 K and 300 mK, undispersed PAni is on the insulator side of the metal–insulator (MI)
. + . +
−
−
. +
−
. +
.+
−
.+
.+ −
. +
−
− . +
−
−
−
.+
− .+
−
−
.+ −
− . +
−
. +
.+ −
−
.+ −
− .+
− −
.+ . +
− . +
FIGURE 1.14
−
.+
. +
−
−
Simplified model for the anisotropic location of the counterions in the metallic core.
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W (=dln(s)/dln(T ))
10
1
0.1
PAni (40%) – PMMA (60%) PAni (33%) – PMMA (67%) Unblended pani
1
10
T (K)
100
FIGURE 1.15 Log–Log plot of W(T) versus temperature for PAni–PMMA blends in the metallic regime. (Reprinted from Wessling, B., Handbook of Nanostructured Materials and Nanotechnology, vol. 5, ed. H.S. Nalwa, Academic Press, New York, 1999, 525. With permission. Copyright 2000 Elsevier Science)
transition. For PAni–PMMA blends, W decreases on decreasing the temperature below 1 K and the systems are found to be on the metallic side of the MI transition. This was the first time that such a material was truly metallic; comparable (CSA-doped) materials were metallic under pressure. In the vicinity of the MI transition, localization and electron–electron interactions both play an important role in determining the conductivity. The results of temperature dependance of paramagnetic susceptibility of PAni and PAni–PMMA blends between room temperature and 2 K further confirm our study. The results are discussed in terms of weak electron–electron interactions in a disordered metal near the MI transition (Figure 1.15). Our commercially available polyaniline powder, ORMECON powder, was used for preparing the PAni–PMMA blends. The magnetic susceptibility of unblended PAni and PAni–PMMA blends was measured using a commercial SQUID magnetometer in the temperature range of 2.0–300 K in an applied magnetic field of 100 mT. The total magnetic susceptibility, xtot, is expressed as a sum of the core diamagnetic susceptibility, xCore, and the paramagnetic susceptibility, xPara: xtot ¼ xCore þ xPara
(1:1)
The dopant is p-toluene sulfonic acid (p-TsA), and the core value of PAni–p-TsA (y ¼ 0.5) is calculated to be 206 106 emu=mol-2ring. The core susceptibilities of PAni and PAni–PMMA blends are calculated by using the core value of PMMA, which is 62.82 106 emu=mol. After subtracting the core value from the experimental values, the total paramagnetic susceptibility of PAni and its blends is plotted as a function of temperature. The data are shown in Figure 1.15 and Figure 1.16. All samples show a nearly temperature-independent magnetic susceptibility down to 50 K. Below 50 K, a temperature-dependent Curie-like susceptibility is observed. Figure 1.16 shows the x vs. 1=T plots for unblended PAni and PAni–PMMA blends. The temperature-independent Pauli susceptibility is calculated from the above plot, and the density of the states at the Fermi energy is calculated. The values of the density of states at the fermi level of PAni (Ormecon) and its blends are much higher than that of all other reported for PAni systems [16,25,57,62], that is, in the range of 2–70 1020, whereby PAni-CSA [27] shows only 0.7–4 emu=eV=2-ring.
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Conductive Polymers as Organic Nanometals
0.0016 #1 #6 #7 #8 fit
c (emu/mole - 2 rings)
0.0014 0.0012 0.0010 0.0008 0.0006 0.0004 0.0002 0.0000 0.0
0.1
0.2
0.3
0.4
0.5
1/T (1/K)
FIGURE 1.16 Pauli susceptibility. (Reprinted from Wessling, B., Handbook of Nanostructured Materials and Nanotechnology, vol. 5, ed. H.S. Nalwa, Academic Press, New York, 1999, 525. With permission. Copyright 2000 Elsevier Science)
The conclusion is clear: without any secondary doping and just by dispersion, PAni–PMMA blends show a temperature-independent Pauli susceptibility down to 50 K, and below 50 K a Curie-like behavior is observed. A finite density of states at the Fermi energy implies that PAni and its blends form a Fermi glass. The Curie-like behavior arises from the single occupancy of localized states at the Fermi energy. We conclude that without any solvent, recrystallization, and secondary doping, dispersion alone, if done properly, can improve the conductivity significantly and enhances the paramagnetic susceptibility in these systems. The results can also be understood with the understanding that PAni as an organic nanometal, is a metal with quantum size-limited metallic space, where the electron wave eventually tunnels from particle to particle. However, these results do not yet explain why PAni in the dispersion processed blend crossed the MI transition to the metallic side, as both raw material on the insulator side, and dispersed product on the metallic side have a much higher density of states then any other comparable material. Obviously, this is only one precondition. Several years ago we proposed a first structure model of the metallic core [28]. For this purpose, we examined seven different PAni powders that showed different conductivities and different dispersabilities. They were all prepared by principally the same procedure (water solution of aniline–p-toluenesulfonate, which is oxidized by potassium peroxodisulfate), with identical monomer–oxidant ratios, but slightly differing reaction parameters.1 Under the assumption of a monoclinic cell, which gave the best fit of the data, a calculation of representative spectra lead to lattice plane distances and indexes of the reflexes. A recalculation of the spectra for the different lots showed slightly different d-spacings (original data in Ref. [29]).
1 The different lots were taken from an international computer-controlled study to evaluate the process below of our polymerisation procedure and to make sure that every lot will be reproducibly the same. We have achieved this necessary reproducibility, which is following the parameters of lot 604, and we are using a computer-controlled reactor with no default results.
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It was very interesting, if not surprising, that the cell volume decreases with increasing conductivity, mainly due to a tendency of decreasing c, which is due to a decreasing angle d of C6-N-C6. To our surprise, it is not the flatter, but the more bent arrangement of the chain that makes PAni more conductive. It seems that with this conformational change, the packing of the chains leads to a better orbital overlap for forming conduction bands. With a C6-N-C6 angle of <1408, it crosses the border of MI transition and becomes truly metallic [24d]. Moreover, the higher order (greater size) of the crystalline regions plays an important role, as can be seen from the evaluation of the half-width of the reflections. These results again showed that the dramatic changes influencing the crystalline structure and then transforming the semimetal into a true metal can take place via dispersion. 1.1.1.3 An Attempt to Determine the Intrinsic Conductivity of Polyaniline on the Nanoscale [30] Electronic and magnetic properties of the emeraldine salt form of polyaniline (PAni-ES) were intensively studied during the last years [2]. There are some conceptions on the charge transport mechanism in PAni-ES. For instance, Houze and Nechtschein [31] have claimed the existence of single conductive chains in PAni, which transport the charges by mobile polarons. Wang et al. [32] have proposed the existence of metal-like crystalline domains characterized as a Fermi glass, in which the electronic states near the Fermi energy, «F, are exponentially localized because of the disorder. The charge is transferred by the phonon-assisted Mott variable range hopping (VRH) [33] between these states. So far, there is no commonly accepted model for the fundamental electronic properties of conducting polymers. As the charge in conducting polymers is transferred by polarons with spin S ¼ 1=2, electron paramagnetic resonance (EPR) spectroscopy is intensively used in the study of conducting polymers [34]. Mobile polarons accelerate electron relaxation of whole spin ensemble. Therefore, the method in principle allows determining the dynamics parameters of such charge carriers directly. An additional change of electron relaxation appears as a result of an interaction of polarons with oxygen molecules possessing sum spin, S ¼ 1 [31]. However, at registration frequencies not more than 10 GHz the method encounters some limitations due mainly to low spectral resolution and strong interaction between paramagnetic centers (PC). These limitations are reduced at the increase of registration frequency up to 140 GHz, so then it becomes possible to identify the nature and transfer mechanism of the charge in various conducting polymers [35]. In the following results of the investigation at wide waveband EPR of magnetic, relaxation, and electronic transport properties of the crystalline phase in PAni heavily doped by p-toluenesulfonic acid (PAni-TSA) are shown. The charge transport was shown [36] to be governed by quasi-three-dimensional (Q3D) charge hopping between mesoscopic crystalline domains with 8 nm size and Q3D extended electron wave functions. Spin relaxation and charge transfer processes were analyzed to be noncorrelated. This contradicts the single conducting chain model and confirms the formation of the Q3D metallike domains in heavily doped PAni-TSA. Commercially available powder like Ormecon polyaniline [37] with y ¼ 0.5 doping level and 30% crystalline fraction was used to study the intrinsic conductivity of PAni on the nanoscale [38]. A quartz capillary was filled with PAni-TSA mixed with diamagnetic MnO powder (1:3) to reduce interparticle interaction. The traces of Mn2þ ions in MnO with geff ¼ 2.00102 and a ¼ 87.4 Gauss were used for the determination of g-factor. At both wavebands, EPR PAni-TSA demonstrates Lorentzian spectrum with Dyson contribution due to the appearance of skin-layer on its surface (Figure 1.17). The temperature dependences of the linewidth of the vacuum-processed PAni-TSA sample determined at both 3 cm and 2 mm wavebands EPR are also presented in Figure 1.17. The data obtained are evidence of the weak influence of the temperature on the linewidth. The diffusion of air molecules into the sample leads to reversible broadening of its EPR line and also to the extremal temperature dependency of the DBpp value with characteristic temperature Tc 160 K
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Conductive Polymers as Organic Nanometals
PAni-TSA
9.7 GHz
140 GHz
20 Gauss
30 3.7 3.6 ∆Bpp (Gauss)
25
∆Bpp (Gauss)
20
2.2
V&G 9.7 GHz 140 GHz
2.0 1.8 1.6
15
100 150 200 250 300 Temperature (K) 10
V&G
air 9.7 GHz 140 GHz
5
0 100
150
200
250
300
T (K)
FIGURE 1.17 EPR spectra of the vacuum-processed PAni-TSA sample exposed to air (solid and dashed lines in the insert, respectively) and the temperature dependence of its absorption peak-to-peak linewidth registered at 3 cm and 2 nm wavebands EPR. The dependences calculated from Equation 1.7 with vdiff ¼ 9.6 1017 exp(0.058 eV=kB T), J ¼ 0.29 eV (dashed line) and vdiff ¼ 3.21018 exp(0.089 eV=kB T ), J ¼ 0.21 eV (dash-dotted line), and C ¼ 0.005 [31] are shown as well. (Reprinted from Wessling, B., Handbook of Nanostructured Materials and Nanotechnology, vol. 5, ed. H.S. Nalwa, Academic Press, New York, 1999, 525. With permission.)
(3 cm waveband EPR) and Tc 130 K (2 cm waveband EPR) (Figure 1.17). The interaction of polarons with oxygen paramagnetic molecules should lead to the EPR line broadening. Figure 1.20 shows the temperature dependences of sac from Dysonian 3 cm and 2 mm EPR spectra. RT intrinsic conductivity of the vacuum-processed PAni–TSA sample lies near 180 S=cm at registration
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4000
vac
air 9.7 GHz 140 GHz
sac (S/cm)
3000
2000
1000
0 100
150
200
250
300
T (K)
FIGURE 1.18 The temperature dependence of ac conductivity, sac, of PAni-TSA determined from its 3-cm and 2-mm wavebands EPR spectra at the absence and presence of air in the polymer using Equation 1.6. The dependences calculated from Equation 1.9 with s0 ¼ 13.2 S cm1K1 and s0 ¼ 6.1 S cm1K1 are shown by dashed and dotted lines as well. (Reprinted from Wessling, B., Handbook of Nanostructured Materials and Nanotechnology, vol. 5, ed. H.S. Nalwa, Academic Press, New York, 1999, 525. With permission.)
frequency of 140 GHz and near to 1500 S=cm at registration frequency 9.7 GHz (Figure 1.18). These values increase up to 1200 and 4000 S=cm, respectively, when the polymer is exposed to air. It should be noted that according to the approach of Q1D polaron diffusion along the single conducting chain [31,39], one should expect the temperature correlation of the PAni–TSA linewidth and intrinsic conductivity. From the data obtained it is seen, however, that the linewidth demonstrates extremal temperature behavior, whereas the sac value changes monotonically in all temperature regions. Therefore, we can conclude that the conductivity of the PAni–TSA sample—as in the case of other PAni samples [35]—is mainly determined by the mobility of electrons Q3D delocalized inside metal-like domains, in which the paramagnetic polarons seem localized on parallel strongly interacted polymer chains. So then, the extreme behavior of temperature dependence of the PAni-TSA linewidth can be explained by the reversible dipole–dipole interaction of localized PC with the oxygen molecules diffusing 0 exp(Ea=kBT ) and an activation energy into the metal-like domains with the frequency of vhop ¼ vdiff Ea. Indeed, it is seen from Figure 1.18 that the dependence of linewidth versus temperature of the PAniTSA exposed to air is well fitted by the following equation: " # hvhop 2 (1:2) d(Dv) ¼ 16=27vhop C 1 þ 12J where vhop is the frequency of the polaron intrachain hopping, C is the number of oxygen molecules per each aniline ring, h is the Planck’s constant, and J is the constant of the spin dipole–dipole interaction. The decrease in Ea value at the increase of the registration frequency should evidence, e.g., for an effect of an external magnetic field on the spin exchange process in the polymer. The shift of the DBpp(T) from extreme to lower temperatures at the increase of the external magnetic field (Figure 1.17) confirms this supposition. The analysis shows that strong and weak spin–spin interaction is realized at T Tc and T Tc, respectively. Also one can conclude that an intrinsic conductivity of the PAni-TSA sample exposed to air can be explained by the Q3D Mott VRH [33] when
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sac (T ) ¼ s0 T
(1:3)
Indeed, Figure 1.18 shows that sac(T) is fitted well by Equation 1.3 with s0 ¼ 13.2 S=cm K1 determined from the analysis of the EPR spectra. The spin–lattice relaxation time of PC in PAni-TSA can be determined by using the CW saturation method as in Ref. [40]. Assuming that polarons diffuse along and between polymer chains with the diffusion coefficients D1D and D3D, one can calculate these values according to literature procedure [41]. RT intrachain and interchain conductivity of the sample were calculated to be respectively 29 and 2.4 103 S=cm. We can conclude that PAni-TSA consists of Q3D metal-like domains with high crystallinity and conductivity embedded into its amorphous phase. The polarons are localized in such domains because of the strong interaction between neighboring conducting chains in crystallites. In contrast to other polymers, this interaction in PAni-TSA is not suppressed by a strong (ca. 5 T) magnetic field. High intrinsic-conductivity of the domains leads to the appearance of the Dyson contribution in their spectra allowing their conductivity to be determined. The intrinsic conductivity is determined by the Mott Q3D intradomain charge transfer. Reversible dipole–dipole interaction of the polarons and oxygen molecules accelerates spin relaxation and increases the intrinsic conductivity.
1.1.2 Nanotechnology with Nanometals Although the above-mentioned results of the investigation of the properties and structures of organic nanometals have been published, and although dispersion as a macroscopic process has dramatic and reproducible effects on the nanoscale, nanotechnological operations and structures have not yet been studied with these nanometals. Research groups active in nanotechnology should begin to look at this new class of materials and, particularly, the dispersion techniques successfully applied to create nanostructures with them. The possibilities with such procedures might be envisioned when looking at the nanoscale structures in dispersion or deposited on various substrates from dispersion. The deposits of the organic metal from a binder-free dispersion (dispersion of the pure raw powder in pure solvents) result in a highly structured layer system: the originally 10 nm primary particles assemble into 50–100 nm big particles, which are laid down in an almost perfect cubic monoclinic lattice (Figure 1.19). It is not totally
+ − + − + −
+ − + − + −
δ+ δ+ Schicht Pet1 06305 300nn
(a)
(b)
+
− +
− +
− − +
+ − +
− − +
− +
− +
(c)
FIGURE 1.19 Particulate substructure of a pure organic metal layer deposited from dispersion. (Reprinted from Wessling, B., Hiesgen, R., and Meissner, D., Acta. Polym., 44, 132, 1993. With permission. Copyright 1993 Wiley– VCH Verlag GmbH)
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Conjugated Polymers: Processing and Applications
unreasonable that precise nanostructures with high variety (but in each case in reproducible form and function) may be prepared, also on an industrial scale, by special deposition techniques using the selforganization ability of the organic metal. In the next section, the deciding basic requirement for the existence and character and for the nanoscale properties of these materials—principal insolubility—will be discussed. This discussion is necessary because the vast majority of the scientific community dealing with the conductive polymers still hopes of obtaining truly soluble conductive polymers’’ one day—a goal that has no chemical or physical basis and would, if it were realistic, prevent their use in nanotechnology. In the further section, based on nanoproperties and nanoscale effects on the macroscale, the applications will be presented and discussed, together with some more visionary potential applications, directed more toward pure nanotechnology.
1.2
Conductive Polymers–Solvent Systems: Solutions or Dispersions
From the beginning of their history in the late 1970s, conductive polymers (organic metals) have been considered as intractable and insoluble. It was an important goal in basic research as in applicationoriented materials science to develop techniques by which they could be processed. The use of solvents was one of the options. As early as 1983–84, after five years of research, we happened to create the first clear dispersions of polyacetylene, polypyrrole, and polyaniline [42], with and without the presence of conventional polymeric binders. This was the beginning of nanotechnology with organic metals. Since 1986, there has been no International Conference on Science and Technology of Synthetic Metals in which the question—solution or dispersion?—did not raise exciting discussions. Most of the scientists support the position [43] that clear (colored) mixtures of (organic) solvents with intrinsically conductive polymers (OM) are solutions. (The first reports of soluble and moldable conducting polymers were published by Elsenbaumer et al. [44a,b].) Only a few support our position [42b], which considers the mixtures to be dispersions. In the recent years, the debate concerning polyaniline and solvent systems has attracted growing interest. During the past 20 years, many publications (in the beginning mainly by the author, later also by other research groups) were published discussing the solubility and dispersibility of the conductive polymers. Only a very short overview of this topic can be given in this handbook. The interested reader can find more detailed explanations in Refs. [15a,15b,45]. Most of the publications proved—even if the authors sometimes made different interpretations of their experimental results—that conductive polymers cannot be made soluble, therefore can only be processed by dispersion. Besides such applications where the polymerization is taking place directly on the substrate and in form of its later use so that no further processing is necessary (like in capacitors), any processing of conductive polymers known to date is based on dispersion. This is the case for all commercially available products, like those from Ormecon (polyaniline in all forms and media from water to organics, polar and unpolar, dispersible or dispersed ready to be used), Bayer (PEDOT, water dispersion), and Panipol (polyaniline, powder dispersible in polyolefines). Even the transformation of an ICP to a true metal (organic metal, OM) proceeds via dispersion (insulator-to-metal transition) [46]. Here, the original conductivity (pressed powder pellet) after polymerization, usually around 5 S=cm, is increased by a factor of at least 10, although the concentration in these predispersion blends is only 40%. The chemical composition is not changed; the dispersion step alone is responsible for the increase of conductivity and the change in the electron transport character that becomes fully metallic (apart from the fact that this organic metal is a nanometal or mesoscopic metal as well, which involves a tunneling contribution to the whole transport). Although several hypotheses have been published in past years [26,46,47], it is not at all clear as to which chain structure and particle morphology are responsible for metallic and high conductivity (and on the other side: nonmetallic and low conductivity) in ICP. None of the many crystal structures published for polyaniline and other ICP seem to appropriately reflect the correct chain structure and
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also have not been helpful in guiding experiments toward a controlled chain alignment that resulted in increased conductivity. Experimental attempts aimed at preparing true solutions of ICP or OM have failed so far. There is no publication reporting about the preparation of single crystals of an ICP or an OM from solution; single crystals have not been possible at all, not by other means as well. The preparation of solutions has been shown to be generally impossible due to thermodynamic reasons; ICP and OM are principally insoluble, according to Ref. [48]. The main reasons for this fact are the unmoldability (no melting point has been found for any of the ICP or OM) which causes the free energy of dissolution to be always >0, due to DG ¼ H(melt) þ H(soln) TS > 0
(1:4)
if DH ðmeltÞ ! 1 the extremely high surface tension (with >150 103 mN=m far beyond any solvent including water) and the very high lattice energy (calculated with a Born–Haber cycle). Therefore, processing of ICP and OM can only be performed using dispersions, whereby the matrix or medium in which the ICP would have to be dispersed, can be a polymer (thermoplastic or thermoset), prepolymer, water, or organic solvent, and the technology has been improved in the last years with the effect that practically all kinds of dispersions in all kinds of media are accessible [2].
1.3
Applications of Organic Metals Emerging from Basic Science: Macroeffects of Nanoparticles
1.3.1 Introduction Practical applications of organic metals (or conductive polymers) like polyaniline (PAni) have been thought of since the very beginning of the research in this field. But what had originally begun with a kind of euphoria, reaching even daily newspapers all over the world, has gone through a deep valley of disappointments since 1989. The strongest motivation for the previously broad support of basic and applied research by the industry and by public funding was the hope of developing polymeric accumulators and batteries using OMs. As this hope vanished (because of the, at best, only equivalent charge storage capacity compared to conventional inorganic systems), most of the bigger companies stopped their work. Earlier attempts by Bridgestone–Seiko in 1987 of marketing a battery [49] (see Refs. [10,32]) were stopped in 1992. Proposed more than 10 years ago, a new concept is attracting a lot of interest, the plastic LEDs and plastic lasers [50], which are based on neutral, not conductive conjugated polymers, where conductive polymers are intended to be used as the hole injection layer. But this area has by far not yet reached a market acceptance as has been hoped for or predicted by the active scientists, companies, and the interested public. Other applications, although at least as challenging and interesting as batteries or LEDs and even much further advanced in the market, have not yet found any comparable attention by the science community, nor the broader public. So, most of our dispersion-oriented research—both fundamentally as applied-oriented—is still widely unknown. No application would ever be realized if OMs could not be processed from an appropriate polymerization over intermediate process steps up to the final product. This problem was far from being trivial, as OMs are insoluble and unmoldable. Moreover, this problem was also not scientifically unproductive: as we managed to find out why they are insoluble and unmoldable, we learned a lot about the important basic properties of the OMs, and we even could overcome these drawbacks, or, as I would say, have learnt to live with insolubility and unmoldability with the help of dispersion. Both areas, batteries and the LEDs, require the OM (or undoped conjugated polymer, respectively) in a pure unblended form as the active material. For batteries, the idea was to polymerize directly into the form of the later electrode. No processing research was thought to be necessary. This was the reason why
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most scientists active in the OM research did not realize the importance of basic research devoted to materials science and processing aspects. With the end of the battery research, the need for processing became evident, as new ideas connected with LEDs emerged.2 These demanded a kind of solvent-based processing technique. On the basis of work by Elsenbaumer et al. [44], a variety of soluble conductive polymers and solvents, even for doped polyaniline, have been proposed. This question was discussed shortly in Section 1.2. Here, only conductive polymers will be discussed. Most other applications or potential applications require a raw material form that can be processed further to the end product. In the meantime, our very early concepts [51] of preparing polymer blends between OMs and various insulating polymeric matrices have been followed by several groups, often using a somewhat different technical approach, but without changing the basic concept: blends are always dispersions of the OM in a matrix polymer, and conductivity above a certain critical volume concentration is possible only because of very complex self-organized structures [15,52]. There are several techniques available for preparing polymer blends (often called composites, which is not the correct term) composed of OMs and an insulating polymer matrix. 1.3.1.1 OM Dispersions and Polymer Blends: Comparative Preparations and Benefits 1. Direct dispersion of a fully polymerized, washed, and dried OM powder (e.g., polyacetylene, polypyrrole, polyaniline, etc., today exclusively practiced with PAni) in a polymer matrix. The first realization of this concept was published by us in 1984 [51]; improvements based on a new technology for the polymerization of dispersible OM powders were later realized in 1987 [53]. The technology described in Ref. [51] and Ref. [53] and in further improvement patents is independent of whether one is preparing a nanoparticulate dispersion in a polymer matrix or in a solvent, or with the help of solvents in a polymer matrix. 2. The technique most widely used in university laboratories is (or has been) the polymerization in situ in the matrix polymer (cf as one of the earliest examples, polyacetylene in LDPE [54]). The disadvantage is that the monomer must be able to diffuse into the matrix, and so must the polymerization agent. Furthermore, there is no possibility of purifying the resulting product, either the resulting polymer, or the blend. Finally, such blends are not processable afterward without losing conductivity. 3. Another (better) option is to polymerize the OM directly on a latex (in fact, S. Jasne from Polaroid had proposed this route in the mid 1980s [55], finally without practical success, and it was again recently proposed by the Intch Company DSM [56]), or in a sterically stabilized colloidal form [57]. Both concepts are based on the idea that an OM polymer blend should be a dispersed system, but that unquestionable idea could get around the very complicate dispersion task by starting with colloidal particles. This is not basically wrong; however, it was not taken into account that
(a) conductivity in a blend is not only a question of the presence of a colloidally dispersed conductive phase, but also of its interfacial structure; (b) the dispersed phase has to have the capability of self-organizing to flocculates, which is only possible with a very specific mechanism,3 first described in Ref. [52] (cf also Ref. [58]); (c) the restrictions of the usability in various polymer blends are caused by the colloidal matrix system (most lattices are waterborne—which might be an advantage if waterbased blends are the only goal—hence, most other water-free matrix polymers are not accessible); and
2 Other product concepts like electrolytic capacitors, as have been realised with TNCQ salts, PPy and PEDT are often approached by polymerizing pyrrole or EDT monomer directly on the substrate. 3 The declamination of the adsorbed matrix polymer monolayer and the formation of a joint layer surrounding all flocculating particles in the complex network structure.
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(d) there were problems with the recovery of the latex or (even more) the sterically stabilized colloidal system before the next process step was begun. We therefore had decided very early before the first proposals of this kind became public, not to follow this route, but to strictly polymerize and recover a dry (dispersible) OM powder, even if the dispersion process itself proved to be the toughest technical and scientific question of all in this whole context. 4. A variation of this concept is to disperse (which is the precise term, in contrast to dissolve) the OM (PAni) in a solvent, preferably with the help of ultrasound to create fully dispersed nanoparticles, and to mix this with a dissolved matrix polymer. Such an approach was described by us in early dispersion patents. Other groups have confirmed the principal feasibility of this approach by counterion induced processability [59], or by dispersing the neutral emeraldine base in a suitable solvent like NMP and mixing this dispersion with a solution of the matrix polymer (e.g., PVP, PMMA, etc.), leading to blends, which had to be postdoped to show conductivity (cf Ref. [60] as one of the more recent examples). It should be noted that in contrast to the opinion expressed by these authors, such blends should never be considered compatible, as the phase size of PAni will be around 50–100 nm, which is not resolvable by light microscope. This is confirmed in Ref. [59b], in which the authors show comparable network structures of aggregated submicron particles, as we have shown earlier [58] for blends resulting from dry dispersion techniques. This also means that solvent (solution) made and processed blends are in fact dispersions. Such blends do not principally differ from other OM blends that are considered nonequilibrium two-phase systems [15], in which the conductive phase is the dispersed (and above the critical concentration flocculated to a dissipative structure network) phase. This has also been supported by the principally equal conductivity and transport properties, as can be seen by comparing the results in Ref. [59b] and Refs. [22b,23a]. 5. A further variation was proposed by the Finnish company Neste Oy and is at the stage of being tested in various ESD applications by Panipol, which attempted to prepare a melt processable polyaniline [61]. It still remains a matter of debate what the melt behavior they observed resulted from, but it was evident that the resulting blend again is a two-phase system with nanosize network structures formed by the dispersed PAni phase. Until now, none of the alternative approaches 2–5 (for a review cf Ref. [62b], a review of processing techniques for conductive polymers) found any practical application except for some minor acceptance of Panipol’s products, and they do not seem to offer an advantage over the dispersion concept favored by us. Although our own research has outlined a complete new theoretical concept, there is still a great need to invest further research into the fundamentals of blend technology, such as dispersion, interfacial phenomena, conductivity breakthrough at the critical concentration, electron transport phenomena in blends, and others. It is not the purpose of this section to review these aspects in greater depth than in Section 1.1 and Section 1.2. In the context of this handbook, it should be sufficient to summarize the basis of any successful OM (PAni) blend with another (insulating and moldable or otherwise processable) polymer is a dispersion of OM (here PAni, which is present as the dispersed phase) and a complex dissipative structure formation under nonequilibrium thermodynamic conditions (for an overview, see Ref. [50]; for the thermodynamic theory itself, see Ref. [15], for detailed discussions, cf Refs. [63,64]). Dispersion itself leads to the drastic insulator-to-metal transition by changing the ‘‘crystal’’ structure in the nanoparticles (see Section 1.1). It is surprising that dispersion generated polymer blends with PAni show typical metallic behavior (see Refs. [20,22b,23a,51]), and some even exhibit an increasing conductivity with decreasing temperature (for the first 50–70 K below room temperature) (see Refs. [22b,23a]) in contrast to the raw PAni used. At any temperature, conductivity is several orders higher compared to the raw PAni (measured as pressed plate). But our dispersion theory is able to explain this phenomenon [62, pp. 566–567]. The new
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Conjugated Polymers: Processing and Applications
findings (Ph–N–Ph angle <1408 for metallic PAni) are a further support for the theses. The macroscopic dispersion step is a tool for very reproducible and crucial processes at the nanoscale. If properly understood, these are the basis for important macroscopic applications based on nanostructures. Many proposals have been made in which OM (mainly PAni) blends could be used. Some of them are visionary and creative, like roofs coated with photovoltaic cells, wallpaper with electrical heating capability, heated textiles, dust filters, and many more [62]. Often the expectation that such blends would have properties superior to those of carbon black–filled blends, in conductivity or in mechanical or colour aspects, guided the vision. Whereas PAni blends can actually deliver somewhat higher conductivity values (up to 50 S=cm, the best value for laboratory samples, see Ref. [22b] and Ref. [23a], 5 S=cm for technical scale [65]) compared to those of carbon black compounds (best values around 0.5 S=cm),4 the other presumed advantages are not there. Nor are mechanical or processing properties, electrochemical stability under applied voltage and current (like for heating devices), or the color aspects of PAni blend any better than with carbon black compounds. There is also often a misunderstanding in the scientific community that carbon black compounds are a relatively bad compromise. This is not the case, as many high performing compounds have been developed and have been in commercial use for many years (cf Ref. [63]). Both the systems, the PAni blends and the carbon black compounds, however, are based on the same structural principle: the conductive phase is the dispersed phase in nanoscale, which suddenly selforganizes into complex networks above the critical concentration [58]. This is the reason for all properties, including mechanical or rheological, and for abrasion. But, most of the poorer products have been replaced by products based on subtle and successful developmental work (cf Ref. [63]). As a consequence, the market does not ask for replacements of carbon black compounds, which provide more or less comparable properties and this at a higher price. (Neste Oy offered PP and PE blends with PAni as a replacement for carbon black compounds in 1995, and stopped the program in 1996.) 1.3.1.2 PAni and PAni=Polymer Blends: Some Key Advantages A new material, like OMs, and blends based on them, will only find a way to those markets, in which either a drastic cost reduction (which is not to be expected) or new useful properties or new useful combinations of properties can be offered. PAni and its blends are new in the following respects: .
. .
.
.
Their conductivity is metallic (more precisely comparable with mesoscopic metals or in other words nanometals). In the galvanic series, PAni and its blends are positioned right below silver. Thin layers (3–50 mm) are transparent, although green, with conductivity values between 109 and 101 S=cm. At higher conductivity, only very thin layers (less than 3 mm) are semitransparent. PAni and its blends are electro- and chemochromic (i.e., they change color upon application of a certain voltage or appropriate chemicals); they also change their conductivity together with these chemical changes: green: metallic; blue: neutral base and insulating; colourless: reduced state and insulating. The three accessible oxidation states reproducibly connected with the three colors (cf Figure 1.57) can be accessed in blends and can serve as different states of PAni in catalytic processes.
Applications like corrosion protection by ennobling and passivation, and manufacture of printed circuit board surfaces composed of the organic metal polyaniline (Ormecon) are based exclusively on this unique set of properties and use most of the properties in parallel. Here, this new materials class is able to offer a performance that cannot be matched by any other material or by another conductive polymer. To realize this performance, commercial products are actually being introduced into the market. A review by Miller [66] is still relatively accurate, except for significant technical and market progress in transparent coatings (Section 1.3.2.1), corrosion protection (Section 1.3.2.2), and printed circuit board 4
Zipperling, Kessler, technical data sheet product 400266, now produced and sold by Clariant Masterbatch GmbH for the commercial use as electrode material in the zincbromine battery of PowderCell, Austria.
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production applications (Section 1.3.2.3). More applications like smart windows (Section 1.3.2.5) and others are still under development and demand further improvements of the basic PAni properties.
1.3.2 PAni Dispersions and Blends Applications 1.3.2.1 Transparent Coatings It is not possible using carbon black compounds to achieve any kind of (semi-) transparent and still antistatic or conductive products or coatings. For packaging and handling purposes of electronic products, however, this is important for preventing electrostatic discharges while still maintaining a capacity for optical inspections. The industry has met this demand by vacuum metallization or humidity-dependent antistats. At least in some areas, OM coatings have been shown and in the future will prove even more to be a well-performing alternative. PAni blend coatings are commercially made by an extreme dispersion of PAni powder in suitable coating systems [67]. The particle size is 70 nm, the critical volume concentration, where the sudden conductivity breakthrough occurs, is around 1%. These are the preconditions under which coatings with a layer thickness of between 1 and 20 mm are transparent (with a transparency of up to 95%). The coatings are green, hence they absorb at 350 nm and, beginning with 650 nm, in the near infrared. Depending on product type and coating thickness, conductivity values between 109 and 101 S=cm are achievable (in resistance values: 109 to 102 V=&). Coating systems for spray, dip, drawbar, gravure, and roller coating techniques, as well as for screen or flexo printing on a variety of substrates (such as plastics including polyolefines, glass, paper, cardboard, and others) have been developed and are in practical use. The various products also offer different hardness, ranging from 2B to 4H. UV and EB curable systems are available. Most of the products are actually solvent-borne (mainly isopropanol, butanol, toluene, or xylene based), but water-borne systems are also becoming available. Comparable systems could be designed for more than just two properties (conductivity and transparency). Electrochromic, indicator, and sensor functions could also be envisioned for coatings based on blends. 1.3.2.2 Corrosion Protection About 10 years ago, deBerry [68] found that a reduction in the corrosion rate could be delivered by electrochemical coating of prepassivated steel (in passivating environment) with polyaniline (from aniline monomer). He supposed that the preformed passive state of the metal was maintained by the PAni layer. Troch-Nagels et al. [69], however, concluded that PAni, unlike polypyrrole, after electrochemical deposition under comparable conditions, did not offer any corrosion protection. Starting in 1986, we tried to coat steel that was not prepassivated, under nonelectrochemical conditions, but with a paint containing dispersed polyaniline. We wondered if some kind of corrosion protection—by whatever mechanism—could be created by an interaction between a dispersion paint and a normal metal surface. This would be, in contrast to previous approaches, a nonelectrochemically applied PAni on a non-prepassivated metal surface. In 1987, we achieved the first promising results [70]. Subsequent work (also published in various patents)5 confirmed the previous findings, but did not show an exciting quantum leap in corrosion protection. Moreover, it was hardly reproducible and did not convince any paint manufacturer. It was only in 1992–1993 that we finally found out after an in-depth evaluation of the interactions between various metal surfaces and coatings of polyaniline (applied as pure dispersion or as dispersion paints) that together with a remarkable corrosion potential shift (ennobling) and an iron oxide layer formation (passivation) lead to a significant anticorrosion effect [71]. In a study together with Elsenbaumer et al. [72], we discovered that the corrosion rate was reduced by a factor of up to 10,000. The iron oxide that formed between the metal surface and the polyaniline primer coating was determined to be Fe2O3, later confirmed with even clearer x-ray photoelectron spectra (XPS) [23b]. 5
For example, PCT=US 93=00543, by Zipperling Kessler and Allied Signal, priority January 1992.
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EB 0.5 O2
4H+ ES + 4H+
2 OH−
4e, 2H+
LE
3 H2O
4e 2 Fe
2 Fe2+
2 Fe3+
Fe2O3
4e O2 + 2 H2O
4 OH−
FIGURE 1.20 Reaction scheme. (Reprinted from Wessling, B., Synth. Met., 93, 143, 1998. With permission. Copyright 1998 Elsevier Science)
We have evaluated the reaction mechanism by which polyaniline as a redox catalyst mediates the reaction between iron (or in analogous way, other metals) and oxygen or water to form the passivating oxide layer [73] (see Figure 1.20). In the meantime, parallel product development toward commercially useful and competitive anticorrosion coating systems has led to various products finding their place in the market, having been successfully tested under practical and various laboratory conditions [74]. It also became evident that the conclusions that were drawn from the basic research on dispersion and the corrosion prevention mechanism of polyaniline have led to superior performance compared to other systems, which have been proposed as alternative techniques [75,114a,114b,114c]. This is probably due to the fact that the alternative methods do not fulfill all chemical, physical, and technical requirements that a corrosion prevention technology based on polyaniline on a technical scale has to. We conclude from our basic chemical, physical, and theoretical evaluations of the polyaniline interactions with metal surfaces leading to corrosion prevention, and from our practical experience with the development, testing, and marketing of various PAni-containing anticorrosion paints (primers) in numerous applications, in widely differing corrosion environments, and summarizing the advantages and disadvantages known up to now that the following requirements have to be fulfilled by a PAnicontaining coating for principally successful (and commercially attractive) applications: . .
.
. .
.
.
The paints must contain well-dispersed (maximum 70–100 nm particle size) PAni. They must be conductive (emeraldine base-containing paints have no ennobling or passivation, but a nice inhibition effect, leading to an anticorrosion efficiency many orders of magnitude lower compared to the conductive form) and metallic; hence they must be a nanometal. The PAni coating (regardless if pure or dispersed in paints) must adhere well on the metal surface, especially under the corrosion conditions. The paints must exhibit metallic properties to ennoble the metal surface. They must offer chemical reactivity (catalytical activity for the reaction path) throughout the whole primer volume. The primer and the complete coating system must be chemically and mechanically stable, also with regard to (interfacial) adhesion (Ormecon polyaniline and their Corrpassiv paints are practically infinitely stable; polypyrrole is not stable, even under ambient conditions). Complete systems have to be designed for each application field and technique,6 because a PAni primer itself does not work properly enough alone or under all practical applications.
6 Ormecon develops and produces specially designed coating systems (Corrpassiv) for steel, aluminium, galvanized steel, and other metals; for general industrial or very aggressive (chemical) corrosion environments; for maritime corrosion; and for various application techniques, like brush or spray coating, roller coating, dip and spin coating, coil coating, and others.
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The design and production of PAni-containing coating systems with commercial applicability asks for much more than just the application of PAni onto a metal surface by whatever method. It requires a full and deep understanding of corrosion prevention chemistry and physics itself, of the interfacial interactions between the PAni primer and the metal surface as well as those between the different coating layers, and especially of the science of OM dispersion. In essence, only 1% of OM in the paint, which has to form a 20 mm primer layer containing a nanosized complex OM particle network, provides a drastic change in the metal surface behavior: it stops corroding. This can be measured, for example, by scanning the voltage potential and by impedance spectroscopy [76]. Here, the nanostructures cause the following macroscopic consequences: . .
A reduction in the delamination velocity from 30–60 mm to close to zero (3 mm=h) An increase in the overall coating resistance by many orders of magnitude
These elements are the basis for several exciting applications [77] that have been achieved with polyaniline anticorrosion coatings. One of the products has been designed as a coating for boats and ships (Skippers CORRPASSIV [78]). It consists of three coating layers, a polyaniline dispersion primer, an interprimer, and a topcoat, in different compositions, depending on whether it is to be applied above or below the waterline. In comparison to well-performing coatings of the previous state of the art (barrier coatings), the new coatings offer a lifetime of about five times longer for the coated vessel. Many hundred boat owners, several shipyards, and bigger container vessels are using the product in real life under commercial conditions. A continuous test in an Icelandic harbor shows its superior performance over that of previously leading anticorrosion coatings. Another heavy corrosion environment is the wastewater management. A tow-layer coating comprising a polyaniline dispersion primer and an epoxy topcoat (CORRPASSIV 4900) are being used and specified for suburban hydraulic wastewater management systems, allowing the replacement of stainless steel hydraulic cylinders with ordinary steel cylinders, with an improvement in the corrosion performance. Furthermore, in a hydraulic system being used in airbus manufacture at Airbus Industries, Hamburg, stainless steel has been replaced by general steel coated with the same product. Variations of this product (for outdoor exposure and for application on prerusted surfaces.) have been designed and used in first real applications in biofilters, wastewater treatment facilities, on several industrial sites, bridges, and a pipeline.7 A further advantage of polyaniline dispersion coatings and complete coating systems including interlayers and topcoats is its applicability to metal substrates other than steel. Aluminum is subject to so-called filiform corrosion, which can be effectively beaten by suitable products, in this case CORRPASSIV 4901 (a different kind of primer and top coat); the performance was tested by an independent institute (see footnote 7). In the meantime, the same primer was also successfully tested in various applications on magnesium.
1.3.2.2.1
A Quantitative Method for Corrosion Performance Studies [79]
To overcome the rather qualitative and poorly reproducible corrosion test methods, a new method, which allows quantitative assessment, was developed by us. Body panels (150 80 1.5 mm3) from Mercedes-Benz were used as substrates. The panels were sand blasted and degreased with a 1:1:1 mixture of acetone, ethylacetate, and xylene. The polyaniline containing primer CORRPASSIV was applied using pneumatic spray units, the thickness after drying averaged 28 + 3 mm. The topcoats were sprayed on after the primer had dried for 24 h. The following topcoats were used on the CORRPASSIV primer with a layer thickness after drying given below: 1. Two-component epoxy topcoat—amine hardened ¼ 2-C EP(154 + 7 mm) 2. Two-component acrylic topcoat ¼ 2-C AY(142 + 6 mm) 3. One-component acrylic topcoat ¼ 1-C AY(140 + 10 mm) 7
Corpassiv test report on filiform corrosion by Forschungsinstitut fu¨r Pigmente und Lacke, FPL, Stuttgart, July 1996.
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The two-component epoxy topcoat, 2-C EP, without and with zinc-rich epoxy primer (56 + 5 mm) was applied on the panels in the same way: 4. Two-component epoxy topcoat ¼ 2-C EP(144 + 14 mm) 5. Zinc-rich epoxy primer ¼ Zn-EP and epoxy topcoat 2-C EP(152 + 15 mm) The three two-layer systems 1–3 were subjected to the salt spray test. The panels were scratched with a scratching tool (Sikkens). The back and the edges of the substrates were covered with an adhesive film. The prepared panels were placed in a spray chamber (Erichsen, Model 608=1000 l) and continuously sprayed with 5% NaCl solution (pH 7). During the test, the chamber was heated to 358C. At certain times, the panels were removed and visually assessed for underfilm corrosion (DIN 53 167), blistering (DIN 53 209), and degree of rusting (DIN 53 210). The fast Fourier transformation (FFT) electrochemical impedance spectrometer model EIS-6416b is a first routine instrument according to Ref. [80]. The experimental setup is given in Figure 1.21. A frequency-rich perturbation signal with a small amplitude is applied to the electrochemical cell controlled by an EG&G Princeton Research potentiostat Model 263A. In the present investigation, a computer-programmed sum of 42 sine waves distributed over four decades was used to synthesize the perturbation signal of a homebuilt signal generator. The peak-to-peak amplitude of the perturbation voltage was usually 15 mV. In some cases, perturbations up to 150 mV were used. The perturbation and the response signal are amplified and filtered by a Stanford Research Systems, Inc. Model SR 640 dual channel low-pass filter. A=D and D=A conversion, timing and controlling were carried out by a 16-bit, 100 kHz transient recorder PC-card from United Electronic Industries, Inc. Model Win-30=3016. Impedance spectra were evaluated by FFT of the perturbation and the response signal (Figure 1.22). All measurements with EIS were performed in three-electrode compartments similar to the cell described in Ref. [81]. The coated substrates were transferred into a chamber with a relative humidity of 96% for 2 days before they were used as working electrodes. A Teflon body separated with a viton sealing ring was pressed on the samples to form the cell. The electrodes had an apparent surface area of
I
Amplifier and Filter
Oscilloscope
Out 1
U
In
2
ADC-card PC DAC-card
U
I
Out CE
In Printer
Signal generator Potentiostat
RE
Cell
WE
IEEE 488
FIGURE 1.21 Setup for FFTelectrochemical impedance spectroscopy. (Reprinted from Wessling, B. and Posdorfer, J., Electrochim. Acta., 44, 2139, 1999. With permission. Copyright 1999 Elsevier Science)
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P−P amp 1. =176.3 mV
5.0 Perturbation
ok
2.5 0.0 −2.5 time, s
−5.0 0.000 5.0
0.999
Response
1.999
2.998
3.998
P−P amp 1. =3.2 uA
ok
2.5 0.0 −2.5 time, s
−5.0 0.000
0
0.999
1.999
2.998
3.998
Perturbation, dB
−10 > −20 −30 −40 10 u 0
f, Hz 100 u
1m
10 m
100 m
1
10
100
1k
10 k
100 k
10 m
100 m
1
10
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1k
10 k
100 k
Response, dB
−10 −20 −30 −40 10 u
f, Hz 100 u
1m
FIGURE 1.22 Time domain and power spectra for perturbation and response signal. (Reprinted from Wessling, B. and Posdorfer, J., Electrochim. Acta., 44, 2139, 1999. With permission. Copyright 1999 Elsevier Science)
16.33 cm2. A platinum wire (Chempur, 99.9%) served as counterelectrode and a silver or silver chloride reference electrode (in 3 mol=l KCl solution, Mettler Toledo InLab 301) was used. All measurements were carried out in 5% NaCl solutions at 258C. The probe measured with SKP was a steel panel ST 14 (85 85 1 mm) from Chemetall. The left side of the panel was coated with a conventional epoxy primer and the right side with the CORRPASSIV primer at a thickness of about 20 mm. After drying for 24 h, the panel was topcoated with the two-component epoxy lacquer at a thickness of about 100 mm. After 1 week of curing at room
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temperature, a defect was applied on every side at a distance of 12.5 mm from the boundary between the two different primers. The scratches (1 40 mm) were applied by cutting the coatings with a knife, removing them with the end of a 1 mm broad microspoon, polishing the bare metal surface with silicon carbide paper (grit 1000), and degreasing them with ethanol. The steel sample was then immersed in 5% NaCl solution for 100 h. The measurements were carried out with a scanning Kelvin-probe from UBM Messtechnik GmbH. The aim of this study was to investigate the compatibility between CORRPASSIV primer and topcoats based on various binders [82]. The selected topcoats have subjected to the salt spray test performed on scratched panels for the three coating systems. Underfilm corrosion, blistering, and degree of rusting were assessed. The test showed that all three coating systems provide highly efficient corrosion protection and complete suppression of underfilm corrosion over a period of more than 1000 h; especially with topcoat 2-C EP, no change whatever was found even after 1000 h. This must be regarded as an indication of the good corrosion control properties of the CORRPASSIV primer. In the case of 1-C AY, however, there was a marked tendency to form blisters, and this made itself felt after only 72 h. The same applied to 2-C AY topcoat, though here the blisters were not observed until after 337 h. The results of the salt spray test indicate a general need for individual investigation of the compatibility of the topcoat with the CORRPASSIV primer. The use of electrochemical impedance spectroscopy (EIS) for the assessment of coatings has been under development for many years now [83]. The first period application of the technique was restricted to the study of the corrosion occurring with poorly protective coatings and coating impedances were evaluated by shape of the Nyquist plots. However, relatively few reports have been published describing results obtained from protective organic coating systems as used commercially [84]. This is surprising as the method is particularly well suited for studying the progressive deviation from the purely capacitive behavior exhibited initially by most such coatings when applied to steel substrates. In our study, three different coating systems with an expected range of corrosion protective properties were used. Figure 1.23a through Figure 1.23c show experimental EIS data for the three different topcoated steel panels as a function of immersion time. Changes in the impedance spectra were observed over a 6 d period except for the 1-C acrylic topcoat with a 1 d period. The curves represent data collected between 50 mHz and 1 kHz immediately on immersion and after 24, 48, and 140 h of testing in 5% NaCl. EIS data for polymer coated metals are usually determined at the corrosion potential Ecorr [85]. As Ecorr cannot be measured for very protective coatings, a potential that is close to that for bare metals is reasonable. For our substrates in aerated NaCl, a value of 560 mV versus Ag or AgCl was applied. Under the corrosive conditions, the coatings of 2-C epoxy and 2-C acrylic topcoat behaved similar in the beginning. During an immersion of 6 d, changes occur for the 2-C acrylic-coated sample. The capacity increased and the resistance decreased whereas they did not change for the 2-C epoxy during the test duration of 6 d. In contrast to these coatings, for the 1-C acrylic topcoat, the capacity and the resistance changed dramatically in a duration of only 1 d. The reason for this is a microscopic damage of the coating, the so-called underfilm corrosion. For comparison changes in the systems using no primer or a zinc-rich primer with the same 2-C EP topcoat, the Bode plots are shown in Figure 1.24a and Figure 1.24b. During the first 70 h, the systems seemed to behave similar to the corresponding CORRPASSIV=2-C EP system, but changes occur after 140 h. The results assessed from the impedance data (Nyquist and Bode plots) in Figure 1.23 and Figure 1.24 after an immersion of 1 week are summarized in Table 1.1. Only for the CORRPASSIV=2-C EP system the coating resistance and capacity remains constant. For all other systems, including the 2-C EP topcoat without primer, the resistance decreases by a factor of 10 and the coating capacity rises by a factor of about 2. For the Zn-EP=2-C EP system, the resistance changes by a factor of 100. The changes cannot be explained only by water absorption of the topcoats. Compared with the systems 2-C EP and Zn-EP=2-C
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3.0E+07
2.5E+07
2.0E+07
t=0h
t = 24 h
t = 48 h
t = 140 h
Zre [Ω]
−Zim [Ω]
2.0E+07 1.5E+07
1.0E+07
t=0h 1.0E+07
t = 24 h
t = 48 h
t = 140 h
5.0E+06
0.0E+00 0.0E+00 5.0E+06 1.0E+07 1.5E+07 2.0E+07 2.5E+07
0.0E+00 0.01
0.1
1
Zre [Ω]
(a)
100
1000
Bode plot
Nyquist plot 3.0E+07
2.5E+07
t=0h 2.0E+07
10
Frequency [Hz]
t=0h
t = 24 h
t = 48 h
t = 24 h
t = 48 h
t = 140 h
t = 140 h
Zre [Ω]
−Zim [Ω]
2.0E+07 1.5E+07
1.0E+07
1.0E+07 5.0E+06
0.0E+00 0.0E+00 5.0E+06 1.0E+07 1.5E+07 2.0E+07 2.5E+07
0.0E+00 0.01
0.1
Nyquist plot 2.5E+07
t=0h 2.0E+07
10
100
1000
Bode plot
3.0E+07
t=0h t=1h
1
Frequency [Hz]
Zre [Ω]
(b)
t = 0.5 h
t=1h
t=3h
t=6h
t = 0.5 h t=3h
t=6h
2.0E+07
Zre [Ω]
−Zim [Ω]
1.5E+07
1.0E+07
1.0E+07 5.0E+06
0.0E+00
0.0E+00
0.0E+00 5.0E+06 1.0E+07 1.5E+07 2.0E+07 2.5E+07
(c)
Zre [Ω]
0.01
0.1
1
10
100
1000
Frequency [Hz]
FIGURE 1.23 Nyquist and Bode plots of selected top coats with CORPASSIV primer for different times of immersion in 5% NaCl solution. (a) 2-C epoxy topcoat; (b) 2-C acrylic; and (c) 1-C acrylic topcoat. (Reprinted from Wessling, B. and Posdorfer, J., Electrochim. Acta., 44, 2139, 1999. With permission. Copyright 1999 Elsevier Science)
EP using the same topcoat, the CORRPASSIV=2-C EP system should change as well. Thus, the changes must be assigned to other changing ion concentrations in the coatings, i.e., iron, zinc, and hydroxide ions. According to these findings, a corrosion protective effect can be assigned to the CORRPASSIV
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Bode plot 3.0E+07
t=0h
t = 70 h
t = 140 h
Zre [Ω]
2.0E+07
1.0E+07
0.0E+00 0.01
0.1
(a)
1 10 Frequency [Hz]
100
1000
Bode plot 3.0E+07
t=0h
t = 140 h
t = 70 h
Zre [Ω]
2.0E+07
1.0E+07
0.0E+00 0.01 (b)
0.1
1
10
100
1000
Frequency [Hz]
FIGURE 1.24 Bode plots of coated steel panels for different times of immersion in 5% NaCl solution. (a) 2-C epoxy topcoat without primer and (b) zinc-rich epoxy primer and 2-C epoxy topcoat. (Reprinted from Wessling, B. and Posdorfer, J., Electrochim. Acta., 44, 2139, 1999. With permission. Copyright 1999 Elsevier Science)
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TABLE 1.1 Change of Resistance and Capacity of Coating Systems with Immersion Time Derived from Impedance Data of Figure 1.23 and Figure 1.24 CORRPASSIV= 2-C EP t(h) 0 0.5 1 3 6 24 48 70 140
CORRPASSIV= 2-C AY
CORRPASSIV= 2-C AY
2-C EP
Zn-EPO=2-C EP
R [kV= cm2]
C [pF= cm2]
R [kV= cm2]
C [pF= cm2]
R [kV= cm2]
C [pF= cm2]
R [kV= cm2]
C [pF= cm2]
R [kV= cm2]
C [pF= cm2]
1613 — — — — 1574 1557 — 1549
51 — — — — 53 52 — 53
1426 — — — — 1076 1049 — 220
89 — — — — 150 154 — 169
1397 964 356 204 53 — — — —
46 60 69 356 204 53 — — —
1136 — — — — — — 1081 123
48 — — — — — — 61 88
1392 — — — — — — 960 16
57 — — — — — — 67 99
primer. These results demonstrate that EIS is capable of detecting localized corrosion phenomena such as disbonding of the coating and initiation of corrosion at the metal or coating interface as well as deterioration of the polymer coating and loss of its protective properties in an early stage. For intact coatings, the major deterioration mechanism is the transport of electrolyte solutions to the interface between the steel substrate and the coating. This transport results from the equalization of concentration and temperature differences. Water molecules disrupt the bonding and polar interactions that are responsible for a good adhesion at the interface [86]. Underfilm corrosion is possible because of the loss of adhesion. The coatings have been defined as intact when the potential of the steel panel was not measurable at the beginning of the test. To study the effects of a coating defect on corrosion at the metal or coating interface and on coating delamination, studies are in preparation with epoxy-coated samples where a small hole with a diameter of 0.5 mm is drilled through the coating into the metal. The Volta-potential measured with a Kelvin sensor is suitable for noncontact measurements of corrosion and surface potentials even under undamaged surface coatings [87]. The function principle and experimental setup is shown in Figure 1.25. The measurement object, the working electrode, and the reference electrode of the Kelvin probe form, due to the small gap between them, a capacitor. Between them a potential is developed, the amplitude of which gives a measure of corrosion activity. A periodic variation in separation by means of an actuator built into the sensor changes the capacitance of the setup. The resulting signal is converted to a measurement signal by means of a lock-in amplifier [88]. The Volta-potential difference is directly determined by the corrosion potential [89]. The corrosion reactions taking place in the vicinity of defects of coatings from pretreated steel substrates are analyzed with the scanning Kelvin-probe. The use of a Kelvin-probe allows measuring of the local electrode potential with high accuracy even below highly insulating polymer layers close to the defect, as the technique requires no electrolytic contact between reference and working electrode. At a defect, metal dissolution is possible (active electrode), whereas below the adherent coating, the dissolution reaction is inhibited (passive electrode) [90]. As shown in Figure 1.25b, the coupling between active and passive parts of the same metal surface forms a galvanic cell. It has been shown that the electrochemical coupling of these active and passive sites determine the delamination rate of the polymer [70]. On the left part, the coating of the metal surface is removed and the metal is in direct contact to an electrolyte. A sharp boundary limits this defect. On the right side of the boundary, the polymer adheres well to the metallic substrate. The potential distribution is measured with the scanning Kelvin probe starting at the defect; the atmosphere within the environmental chamber is kept at a relative humidity of 100%.
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Vibrating reference electrode
w
∆ψ
i
Polymer Metal working electrode
(a)
∆U O2
O2
Electrolyte
Me2+
Me
(b)
Local anode
Polymer Oxide
C+ ze−
Metal
Local cathode
FIGURE 1.25 Corrosion potential measurements using a scanning Kelvin-probe. (a) Function principle and (b) galvanic element forming under delaminated metal=coating interface [90]. (Reprinted from Wessling, B. and Posdorfer, J., Electrochim. Acta., 44, 2139, 1999. With permission. Copyright 1999 Elsevier Science)
The potential distribution between two defects on a steel sample was measured. On the left side, the sample was coated with a conventional primer and on the right side with the CORRPASSIV primer, both topcoated with a two-component epoxy lacquer at a thickness of about 100 mm. The transition between negative and positive potentials marks the delamination front. As shown in Figure 1.26 the delamination zone is not an equipotential-surface. More positive potentials are registered toward the delamination front. This is because of the changing reaction conditions within the delaminated area, which depends on the distance to the defect. In the curve of Figure 1.26, the delamination zones are clearly to be seen. The delaminated area on the left side coated with a conventional primer is larger than the zone of area coated with the CORRPASSIV primer by a factor of two. Due to the polyaniline-containing coating, the measured potential on the right side is shifted to more positive potentials. As shown, the combination of the methods EIS, scanning Kelvin-probe, and salt spray test with the addition of a climate cycling test according to VDA (Association of German Automobilists) is a productive tool for testing and performing corrosion protection agents and lacquer systems. In contrast to salt spray and climate cycling tests (2–3 months), useful results from measurements with EIS and the scanning Kelvin probe can be obtained within 1 or 2 weeks, from where real-world performance predictions can be derived. 1.3.2.3 Printed Circuit Boards After 3 years of development, including 2 years of practical testing in the industry up to a commercial scale, we can report (text written in 1998) an interesting new application: the preparation of the surface finish of printed circuit boards. The places where the diodes, resistors, etc., have to be mounted and connected to copper-based printed circuits have to be solderable for a period of at least 1 year in
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E/mV
Delamination zone
Delamination zone
+100
0
−100
−200 2-C EP top coat
Defect
Conventional 2-C EP primer
−300
30
20
10
0
Defect Corrpassiv
Steel
10
30
20
d/mm
FIGURE 1.26 Potential profile as measured with the Kelvin-probe after an immersion time of 100 h in 5% NaCl solution. (Reprinted from Wessling, B. and Posdorfer, J., Electrochim. Acta., 44, 2139, 1999. With permission. Copyright 1999 Elsevier Science)
all climates. This is tested in the industry by >4 h at 1558C, followed by a complex solderability testing procedure. Copper alone does not meet this requirement, whether it is coated with organic coatings (like benzotriazoles) or chemically deposited tin. Only gold (with a Ni interlayer), palladium, or a thick layer (10–17 mm) of solder tin (applied by a melt and then hot air leveling process) do the job. We have found that a very thin (80 nm in average) coating of PAni—deposited from a water dispersion8 followed by a special chemical tin deposition of only 0.5–1 mm thickness—allows us to achieve solderability stability for more than 4 h at 1558C—ageing plus a full industrial soldering process. Here, PAni’s function is fourfold: (a) it passivates the copper (see Figure 1.27 [91] and Figure 1.28); (b) it catalyzes the tin deposition so that also with the increased Cu concentration used tin deposition baths, only pure tin is being deposited (see Figure 1.29); (c) the interdiffusion of Sn and Cu is inhibited; (d) the oxidation of the final tin surface is decreased. This application is based on the most advanced dispersion technology for PAni in water. In contrast to other approaches, nonfunctionalized polyaniline has been used as the raw material. This makes it possible to start with a commercially available polyaniline powder (Ormecon) instead of special polyaniline derivatives for different applications or dispersion media. It shows, furthermore, the generally applicable dispersion concept. The application on printed circuit boards is a long-term and high-end application and is specified on the basis of the long-term stability of our product.
8
This is the first commercial dispersion of PAni in water (cf Ormecon GmbH data sheet Ormecon CSN, PCB 7000).
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Current density [mA/cm2]
Conjugated Polymers: Processing and Applications
20 18 16 14 12 10 8 6 4 2 0 −1.2 −2 −4 −6 −8 −10 −12 −14 −16 −18 −20
Cu uncoated
Cu coated with ORMECON
−0.2 −1
−0.8
−0.6
−0.4
0
0.2
0.4
0.6
0.8
1
1.2
1.4
1.6
1.8
2
2.2
2.4
Potential U−H2 [V]
FIGURE 1.27 Potential curves of copper. (Reprinted from Wessling, B., Synth. Met., 93, 143, 1998. With permission. Copyright 1998 Elsevier Science)
1.3.2.3.1
Oxidation of Copper in the Presence of the Organic Metal Polyaniline [92]
During storage, unprotected copper surfaces are oxidized, leading to the formation of Cu2O and later to a mixture of Cu2O and CuO. By coating copper with organic agents (imidazole, benzotriazole, or PAni), the surface can be protected [93]. When treated with a dispersion of PAni, copper is covered with a thin film and the corrosion rate of the metal is reduced [94]. Upon aging at high temperatures, various colors are developing. We studied the mechanism of PAni–Cu interaction. Results are presented on the formation of Cu(I) complexes, Cu(I) and Cu(II) oxides with polyaniline, and without any surface protection. The copper foil (99.999% Goodfellow) used was polished using 6, 3, and 1 mm diamond paste. It was rinsed with ethanol and cut into 1.5 1.5 cm2 pieces. Copper was etched for 1 min in 1 M sulfuric acid, rinsed with deionized water and (a) air dried or (b) pretreated by immersing in a water dispersion of OM (0.01% and 3% PAni, Ormecon GmbH) for 1 min, rinsed with water and air dried. They were annealed in a convection oven (Nabertherm) at 1558C for different hours. Electrochemical experiments were performed using a potentiostat or galvanostat (EG&G, model 263A) computer controlled by an IEEE-488 GPIB interface board. For chronopotentiometric reduction of copper oxide films, a current of 100 mA=cm2 was applied. The electrochemical reduction was performed in a borate buffer at a pH of 8.5 in a 50 ml glass cell with three 14.5=23 standard tapers. The copper foil served as working electrode, a platinum wire as counter electrode and a Ag or AgCl (3 mol=l KCl) as reference electrode, all mounted with taper joints. Surface morphology and appearance were determined by scanning electron microscopy using a Philips XL30 SEM. Ageing of copper pretreated with the OM dispersion in comparison with sulfuric acid etched plates causes the color to change with annealing time from various violet and bluish tones into silver and bright golden. This is due to interference of light between the upper and lower thin oxide layers. Varying oxide layer thickness causes different colors. The copper oxide thickness was measured by sequential electrochemical reduction [95]. The applied electric current is directly proportional to the oxide mass, provided the reduction current and the depleted area remain constant. A relationship between color and oxide layer thickness was found. The change of oxide layer thickness for Cu2O and CuO is shown in Figure 1.30. The growth of copper oxides on surfaces etched with acids is proportional to time as expected for oxidation of metals in the gas phase. The PAni-pretreated surface behaves differently. The formation of p CuO is proportional to t, whereas Cu2O decreases exponentially to a minimum of 0.6 nm. From the p slope of a CuO thickness versus t plot, the rate constant for the formation of CuO was calculated.
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1. H2SO4/Benzotriazole 2. Heating
15 µm Material untreated
Heating
1. H2SO4/PAni 2. Heating
1. H2SO4/PAni 2. Heating
4. H2SO4/PAni 5. Heating
(a)
(b)
FIGURE 1.28 (See color insert following page 8-22.) Copper passivated by Ormecon. It becomes golden with tempering. (Reprinted from Wessling, B., Synth. Met., 93, 143, 1998. With permission. Copyright 1998 Elsevier Science).
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Cu2+ - Increase and Sn2+ - decrease without ORMECON 0.20
Cu2+ and Sn2+ [mol/I]
0.15
0.10
0.05
Cu[II]-increase Sn[II]-decrease Max.limt for Cu 0.00 0.00
0.10
0.20
0.30
0.40
0.50
0.60
Min.limt for Sn 0.70
0.80
0.90
1.00
Tinned surface area [m2/I] Cu2+ - Increase and Sn2+ - decrease with ORMECON 0.20
Cu2+ and Sn2+ [mol/I]
0.15
0.10
0.05
Cu[II]-increase Sn[II]-decrease Max.limt for Cu 0.00 0.00
0.10
0.20
0.30
0.40
0.50
0.60
Min.limt for Sn 0.70
0.80
0.90
1.00
Tinned surface area [m2/I]
FIGURE 1.29 Linear tin deposit on copper after ORMECON pretreatment. (Reprinted from Wessling, B., Synth. Met., 93, 143, 1998. With permission. Copyright 1998 Elsevier Science)
Rate constants at different temperatures were determined for both sample sets. The slope of an Arrhenius plot of rate constant versus 1=T gives an activation energy of 35 kJ=mol for CuO formation. Without PAni, an activation energy of 50 kJ=mol was calculated. p The T dependence occurs because the formation of CuO is a diffusion limited reaction: copper is oxidized by PAni, the copper ions are transported through the film and form a Cu(I)–PAni complex; by
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4
Cu2O/nm
3 2 1 0 0
100
200
300
400
T, min 60 50
CuO/nm
40 30 20 10 0 0 0.01% PAni
100
200 T, min 3% PAni
300
400 H2SO4
FIGURE 1.30 Formation of Cu oxide layers depending on ageing time. (Reprinted from Posdorfer, J., and Wessling, B., Synth. Met., 119, 636, 2001. With permission. Copyright 2001 Elsevier Science)
diffusion Cu(I) ions are transported to the film–gas interface undergoing partial oxidation to Cu(II) easily because of the lower activation energy. Analysis of the surface morphology and appearance by using SEM shows that an uniform copper oxide layer is formed on the PAni-treated sample (Figure 1.31). The untreated copper is very rough (Figure 1.32). During immersion of Cu in a dispersion of PAni, a Cu(I)–PAni complex forms on the surface, which determines the oxidation of the copper and the passivation of the surface. A Cu2O layer with almost p constant thickness is formed, the growth of the CuO layer is proportional to t of ageing.
1.3.2.3.2 Chemical Interaction at Copper and Polyaniline Interfaces [96] Polyaniline is a polymer with a variety of interesting properties and applications [97]. For example, it has been proposed in several reports that a polyaniline layer can be used to inhibit the corrosion of steel surfaces [98]. One aspect of this work is to investigate the interaction of polyaniline with a copper substrate to check how the oxidation behavior of copper is influenced by the polyaniline. Another aspect of this work is to investigate the chemical interaction and the growth of vapor-deposited copper on the emeraldine salt of polyaniline. The copper substrate was a sputter-cleaned pure copper foil. On exposure of the copper foil to air, polyaniline was deposited onto it using a commercially produced polyaniline dispersion in an organic solvent (Ormecon GmbH). The sample was kept in air at room temperature and after a certain time in air was transferred back into the main chamber for analysis. XPS measurements were performed using
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Acc.V Spot Magn 10.0 kV 3.0 40000x
Det SE
WD 4.7
1 µm
FIGURE 1.31 SEM image of golden copper=PAni surface after 2 h at 1558C. (Reprinted from Posdorfer, J., and Wessling, B., Synth. Met., 119, 636, 2001. With permission. Copyright 2001 Elsevier Science)
an electron spectrometer (VG MK II) equipped with a non-monochromatized Al Ka source (energy, hn ¼ 1486.6 eV) and operated at 340 W. The ejected photoelectrons were energy analyzed using a hemispherical electron analyzer at a pass energy of 20 eV corresponding to an energy resolution of 1.2 eV for Ag 3d5=2 line. The base pressure in the analysis chamber during the XPS analysis was better than 1 1010 mbar. For deposition of copper onto polyaniline, the dispersion was dried on a stainless steel foil. Copper was evaporated from a tungsten crucible. For the TEM measurements, thin polyaniline films were prepared on glass slides. To transfer the polyaniline films onto Ni meshes, the glass slides were dipped
Acc.V Spot Magn 10.0 kV 2.4 40000x
Det SE
WD 5.0
1 µm
FIGURE 1.32 SEM image of copper oxidized by annealing at 1558C for 2 h. (Reprinted from Posdorfer, J., and Wessling, B., Synth. Met., 119, 636, 2001. With permission. Copyright 2001 Elsevier Science)
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into distilled water and the floating polyaniline film was caught by a Ni mesh. Copper was then evaporated onto the dried polyaniline film. TEM micrographs were taken with a Philips CM 30 microscope operated at an accelerating voltage of 200 kV. Commercially available polyaniline dispersion was deposited onto sputter-cleaned copper foil. If the deposition is performed under a steady argon flux and the sample is transferred directly back into the vacuum, the green conductive emeraldine salt remains unchanged. This indicates that in the absence of oxygen, no reaction occurs between the polyaniline and the copper substrate. Upon exposure to oxygen, the green emeraldine salt is transformed within 20 s to the blue nonconductive emeraldine base and at the same time Cu(I) oxide formation is observed. Also, on prolonged exposure of the sample to air Cu(II) oxide is formed (see Figure 1.33). This is different to a sample without polyaniline, where after the same time only Cu(I) oxide is detected (see Figure 1.33). It seems that polyaniline enhances the oxidation of the copper. A polyaniline film was dried at room temperature and then removed from the copper substrate. XPS analysis of the polyaniline revealed that some copper was dissolved from the copper substrate and transferred into the polyaniline. Different amounts of copper were deposited onto the conductive emeraldine salt of polyaniline. Careful XPS measurements showed that there are no chemical changes of the polyaniline or of the tosylate anion. For small copper coverages below 0.1 nm one observes a shift of the copper L3M4,5M4,5 Auger line to lower kinetic energies. This shift seems to be not of chemical origin, but due to the formation of small copper clusters or islands [99]. This is supported by the observation that the Auger line is shifted more and more toward the value expected for copper metal as the clusters grow with increasing copper coverage. It is noted that the presence of copper strongly enhances the degradation of the tosylate counterion on exposure to x-rays. In this degradation reaction, the tosylate anion loses its oxygen and the sulfur gets reduced from a þ6 state to a 2 state. In the presence of copper, a noticeable amount of the degradation product is already detected after only a few minutes during an XPS measurement. A similar effect may also be responsible for chemical changes of the perchlorate anion reported in Ref. [100], rather than a chemical reaction induced directly by the copper.
10 min 300 min without OM
Intensity [arb. units]
30 min Increase of the Cu(II) fraction widening of the peak
60 min Decrease of the main peak
Increase of the Cu(II) fraction 300 min With OM
945
Binding energy [eV]
925
FIGURE 1.33 Cu 2p3=2 photoemission spectra of copper foils with and without polyaniline after prolonged exposure to air. (Reprinted from Ladebusch, H., Strunskus, T., Posdorfer, J., and Wessling, B., Synth. Met., 121, 1317, 2001. With permission. Copyright 2001 Elsevier Science)
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60.00 nm
60.00 nm
60.00 nm
60.00 nm
FIGURE 1.34 TEM micrographs of copper deposits on the green conductive emeraldine salt of polyaniline. Nominal copper coverages are 1, 2, 5, and 10 nm, respectively. The substrate temperature was 608C and the deposition rate was 0.1 nm=min. (Reprinted from Ladebusch, H., Strunskus, T., Posdorfer, J., and Wessling, B., Synth. Met., 121, 1317, 2001 With permission. Copyright 2001 Elsevier Science)
The weak chemical interaction of copper prevents the formation of a closed layer at low copper coverages. Typical TEM micrographs are displayed in Figure 1.34. At a coverage of 1 nm, one observes a layer of closely packed copper islands. Increasing the coverage first leads to a coarsening of the copper islands, but up to 10 nm copper coverage the film is almost completely closed. XPS measurements indicate that a coverage of about 15 nm copper is sufficient to form a closed copper film. Under inert conditions, the conductive emeraldine salt of polyaniline does not react with copper. In air, it is transformed fast (within a minute) to the nonconductive blue emeraldine base and copper gets oxidized to Cu(I). In a slower process (within several hours), Cu(I) is then transformed to Cu(II). It was demonstrated that the presence of polyaniline changes the oxidation behavior of the copper. Copper evaporated onto polyaniline does not chemically react with the polyaniline or the tosylate anion and forms a dense layer of copper islands.
1.3.2.3.3
A Superior Whisker-Reducing Immersion Tin Technology [101]
With the growing global demand for immersion tin, the requirements placed on product and process technology are increasing significantly. Market leading OEMs require surface finishes with
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increased tin thickness, improved solderability at multiple cycles and higher temperatures, enhanced surface flatness, fine pitch compatibility, and—ever more important—limited whisker growth. Immersion tin can meet these requirements and has proven its advantages with market-leading PCB manufacturers. When talking about immersion tin as a final surface finish for PCBs, whisker growth is a topic of concern. Ormecon has spent a significant amount of time and resources investigating the whisker formation phenomenon and its prevention. On the basis of this experience, the company has developed a whisker-reducing immersion tin product: Ormecon CSN FF-W. All chemicals used are lead-free and halogen-free. This product, as well as the already established Ormecon CSN and the recently introduced CSN FF, are based on the organic metal polyaniline, which is responsible for a completely different deposition mechanism and tin layer structure. This is due to the fact that the organic metal not only passivates the copper (and later the tin), but also exclusively forms Cu(I) cations and transfers electrons to the Sn(II) ions for the reduction and deposition, thus acting as a catalyst. Whisker formation is a typical feature of metallic tin. There can be two kinds of whisker phenomena: at the component itself and also at the tin surface finish. Various factors can cause whisker formation. One key factor is internal stress. There are two stress categories. . .
Compressive stress, which drives the whisker growth Tensile stress, which reduces the propensity of whisker growth
Internal stress mainly develops during storage as a result of diffusion processes at the Cu–Sn boundary. With tin plated over copper, compressive stress is built up, because copper diffuses much faster into tin than tin diffuses into copper. With continuous diffusion over time, compressive stress increases in the tin layer. Most often, the presence of tin oxides prevents the stress from being released. However, the tin oxide layer is never perfect (defects are the norm) and when the internal stress becomes high enough, it will break through the defect in the oxide layer and a whisker can form to release the stress (Figure 1.35). The products currently available on the market for whisker reduction are mainly based on reduced copper-into-tin diffusion because of the more compact nature of the tin layer. Also whisker-reducing effects are produced by a silver layer on top of the tin layer. Such products, however, do not satisfy some of the more stringent whisker-reduction specifications. Some manufacturers require no whiskers with a length of more than 20 mm and even less.
Whisker SnOx Sn CuSnx Internal stress Cu
FIGURE 1.35 Diagram illustrating whisker formation process. (Reprinted from Arendt, N., Baron, F., Benz, V., Letterer, M., Merkle, H., Schroeder, S., and Wessling, B., OnBoard Technology, Wise Media, 30–32. With permission. Copyright 2004 OnBoard Technology)
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ORMECON CSN FF/FF-W Immersion tin 2Cu0
2Cu1+ 2e− 5 µm
OM
OMred 2e−
Sn0
Sn2+ SEM photo of virgin deposit Highly crystalline morphology Dense package Smooth surface High resistance against diffusion and oxidation
FIGURE 1.36 Ormecon immersion tin deposit reaction (Reprinted from Arendt. N., Baron, F., Benz, V., Letterer, M., Merkle, H., Schroeder, S., and Wessling, B., OnBoard Technology, Wise Media, 30–32, With permission. Copyright 2004 OnBoard Technology.)
The silver-topped tin layer is relatively expensive, as it requires an additional process step and a significant amount of silver. Therefore, a new immersion tin product and process technology for whisker reduction is necessary to meet all the requirements of the market-leading OEMs and PCB manufacturers. As in Ormecon CSN and Ormecon CSN FF, the predip of the Ormecon CSN FF-W product contains the organic metal in aqueous dispersion, which catalyses the tin deposition, resulting in a very dense tin layer with a big grain size morphology. Figure 1.36 and Figure 1.37 show the difference between the Ormecon immersion tin deposit reactions compared to standard immersion tin reactions.
Deposit’s morphology of traditional immersion tin
0
2Sn2+
2Sn
4e−
2Cu2+ 5 µm
2Cu0
Traditional immersion tin
?
2e− Sn0
2Cu1+ Sn2+
SEM photo of virgin deposit Fine grain crystals Porous structure Rough surface Sensitive to diffusion and oxidation
FIGURE 1.37 Standard immersion tin reactions. (Reprinted from Arendt, N., Baron, F., Benz, V., Letterer, M., Merkle, H., Schroeder, S., and Wessling, B., OnBoard Technology, Wise Media, 30–32. With permission. Copyright 2004 OnBoard Technology.)
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Conductive Polymers as Organic Nanometals
5 µm
5 µm
Pores
FIGURE 1.38 Diagram of pores of small grain and large grain tin finishes. (Reprinted from Arendt, N., Baron, F., Benz, V., Letterer, M., Merkle, H., Schroeder, S., and Wessling, B., OnBoard Technology, Wise Media, 30–32. With permission. Copyright 2004 OnBoard Technology.)
The key benefits of the Ormecon deposit are .
.
.
A large grain size tin layer is thermodynamically more stable than tin with a small grain size and it is less susceptible to recrystallization. It is therefore more difficult to squeeze out a whisker from a large grain size tin. A large grain size tin layer furthermore reduces the speed of copper diffusion, and thus less stress is produced compared to a rapid diffusion of copper into a small grain size tin. The big, flat tin crystals of an Ormecon CSN FF=FF-W finish offer significantly less room (pores) for impurity inclusions compared to a small grain size tin deposit. The small tin grain size compared with the Ormecon large tin grain size contains more pores (Figure 1.38).
However, as documented in Table 1.2, this is sometimes not sufficient. Therefore, Ormecon has completely changed the layer structure by adding an additional whisker-inhibiting ingredient in the predip process step. This can be selected from various metals. In the new product Ormecon CSN FF-W, silver is used. Thereby, a metal alloy nanolayer (10–20 nm) containing Ag and Cu is formed on the copper surface, again induced by the action of the organic metal catalyst. The subsequent tin layer deposition is then carried out in a way that a new sandwich layer structure is setup. There is a precisely defined tin layer on top, followed by a broader intermetallic layer. All layers show a much smoother concentration gradient than in other types of tin surfaces. Such a new sandwich layer construction has proven to have a superior characteristic in all respects. On the basis of secondary ion mass spectrometry (SIMS) (Figure 1.39), it was possible to prove that .
.
.
.
Sn can be found at a much greater depth in the Cu layer (up to 3 mm or more) compared to standard immersion tin products, where Sn can only be found up to about 1.5 mm deep; Cu can be found directly at the surface (in very low concentrations) in contrast to standard products, where Cu is only present starting at a depth of about 0.8–1 mm; Ag is not present in the form of a separate layer, not even below the tin layer (which was deposited after Ag deposition), but everywhere in the tin–copper and more deeper in the copper–tin alloy, with a maximum concentration slightly below the surface; the concentration of all three metals gradually changes with the increasing depth.
The fact that Ag is not present in form of a separate layer eliminates the risk of silver migration over the time. TABLE 1.2
Whisker Formation with Various Immersion Tin Surface Finishes
Product ORMECON CSN FF-W ORMECON CSN FF Product A Product B Product C
Storage Time 7 Weeks Whisker Size
Storage Time 14 Weeks Whisker Size
No whisker few 20–30 mm many >100 mm many 20–30 mm several 20–30 mm
No whisker Not measured Not measured Not measured Not measured
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10
Cu Cu Ag Sn CuruSn x 0,001
rel. Intensität
10−3
Cu
10−3
10−4
Sn
10−5
Probe 2
−2
rel. Intensität
10
ORMECON CSN FF_W
ORMECON CSN FF Probe 1
−2
Sn Cu Ag Sn CuruSn x 0,001
10−4
10−5
Ag 10−6
10−6
Ag 10−7
10−7 0
500
1000
1500
Sputterzeit [S] ~1.5 µm Sn
Cu/Sn
2000
0
500
~3 µm Cu
1000
Sn/Cu sandwich Ag rich
Separate Cu/Sn alloy layer
1500
Sputterzeit [S] ~1.5 µm
Ag poor
2000
~3 µm
Cu/Sn alloy With a small portion of Ag
Sn/Ag/Cu sandwich layer with smooth concentration gradient
FIGURE 1.39 SIMS results showing diffusion of tin, silver, and copper in various sandwich structures. (Reprinted from Arendt, N., Baron, F., Benz, V., Letterer, M., Merkle, H., Schroeder, S., and Wessling, B., OnBoard Technology, Wise Media, 30–32. With permission. Copyright 2004 OnBoard Technology.)
Figure 1.40 shows the difference in the layer structure resulting from the use of standard immersion tin products and the new whisker preventing sandwich layer structure. The concentration gradient of both Cu and Sn is much less steep compared to the previous layer structures. Diffusion is dramatically reduced due to the fact that the concentration of Cu in Sn in the surface regions and of Sn in Cu in the deeper regions is relatively high to begin with. This fact, together with the ennobling effect of Ag and Without whisker-protection Predip
Without whisker-protection Predip SnOx Sn Sn/Cu
Smooth Sn/Ag/Cu “sandwich”
Cu
High amount of copper diffusing into tin ⇒ Three separate layers are formed: Sn Sn/Cu alloy Cu ⇒ Compressive stress accumulates in the layer
Decreased amount of copper diffusing into tin/time and reduced overall diffusion speed ⇒ Smooth Sn/Ag/Cu sandwich ⇒ Decreased velocity of intermetallic layer build-up ⇒ Decreased stress formation
FIGURE 1.40 Layer structure with and without the whisker protection predip step. (Reprinted from Arendt, N., Baron, F., Benz, V., Letterer, M., Merkle, H., Schroeder, S., and Wessling, B., OnBoard Technology, Wise Media, 30–32. With permission. Copyright 2004 OnBoard Technology.)
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the organic metal, is also the reason for the extremely good solderability and ageing behavior of the immersion tin layer. This new sandwich structure effectively reduces and even prevents whisker formation, as shown in Table 1.2, because less stress is induced due to the much lower diffusion of both Cu and Sn. This innovative process fulfills the requirements of several currently implemented OEM whisker-reduction specifications. The product was introduced in 2004 with some high volume PCB manufacturers. The PCB manufacturers choosing to apply Ormecon immersion tin technologies can easily change between the standard immersion tin product Ormecon CSN FF and the high-performance, whisker-reducing immersion tin product Ormecon CSN FF-W. The difference between two such products is only the predip step. Due to the deposition process (using noble organic metal and noble silver), the resulting silver-containing sandwich composition and the compact and big crystal grain size structure, the ageing behavior of this surface finish has proven to be superior. The diffusion rate of Cu into Sn at 1558C is even lower (less than 3.3 nm=s) than in Ormecon CSN and Ormecon CSN FF (around 4.0 nm=s, which is significantly lower than with any other conventional immersion tin technique). Oxidation is also strongly reduced. Consequently, after thermal ageing or multiple solder steps, the solderability is preserved to a superior degree. In particular .
.
.
electrochemical analysis (GCM, SERA) shows remaining pure tin is about 0.2–0.5 mm thick (depending on starting tin layer thickness and ageing conditions); solder tests with a meniscograph show a wetting angle of less than 558 after ageing or multiple reflow; and wave solder tests show complete solderability after thermal ageing or multiple solder steps (reflow).
For the electrochemical tin layer thickness measurements (GCM or SERA), samples with comparable virgin tin layer thickness were made: 0.58 mm for both Ormecon CSN FF and Ormecon FF-W. After thermal ageing at 1558C for 4 h, a tin thickness of 0.17 mm remains on the standard surface, whereas for the new Ormecon FF-W, a thickness of 0.24 mm remains, corresponding to the lower diffusion rate. Figure 1.41 shows the process technology of the two products. For an accurate process control, Ormecon recommends the measurement of the layer thickness with the coulombmetric measurement system (GCM, SERA), before and after ageing. 1.3.2.4 Electro- and Chemochromic Applications and Other Sensors Potential applications based on the different colors of OMs linked with their different oxidation states have been studied for a long time. PAni, for example, is green in its doped metallic salt form, blue in the neutral base form, and colorless in the reduced stage (Figure 1.42). One can switch between these stages by applying a certain voltage or with appropriate chemical agents (acids, bases, reductive or mild oxidative agents, respectively). Electrochromic windows, sensors,9 and indicators could be created. (For an overview, see Ref. [102].) Up to now, however, no significant application-oriented development has taken place, at least as far as being published. Several announcements did not come true. For future development, it might be interesting that some blends, which are used for transparent coatings (cf Section 1.3.2.1) respond to electrochemical switching of oxidation states (and hence color) at least as quick as pure PAni layers (A. Wappner, personal communication). Coatings from blends or from pure dispersions are easily applied and will serve as technically fully satisfying and commercially attractive processing tools. Direct polymerization on ITO glass is actually preferred by R&D groups but will not be reproducible enough and not at all competitive to dispersion or blend coating techniques, because of difficult process control, purification, and chemical waste problems in factories, where chemical processes are usually not practiced. 9
Recently, a first gas sensor has been constructed and principal gas detection (type, concentration) has been demonstrated, using our PAni dispersion by M. Meijerink, University of Neuchaˆtel (personal communication, to be published).
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Dry
Dl water rinse
Immersion tin
Rinse
Predip
Warm rinse
Micro etch
Rinse
Acid cleaner
Rinse
Conjugated Polymers: Processing and Applications
ORMECON CSN FF and ORMECON CSN FF-W consist of 4 active steps: Acid cleaner
ACL 7001
Micro etch
MET 7000
Predip (Organic metal based)
OMP 7000 OMP 7001
Immersion tin
CSN 7004
For ORMECON CSN FF For ORMECON CSN FF-W
FIGURE 1.41 Process technology for CSN FF and CSN FF-W. (Reprinted from Arendt, N., Baron, F., Benz, V., Letterer, M., Merkle, H., Schroeder, S., and Wessling, B., OnBoard Technology, Wise Media, 30–32. With permission. Copyright 2004 OnBoard Technology.)
1.3.2.4.1
Electrical Conduction in Polyaniline–PMMA Blends and Their Use as Cryomagnetic Temperature Sensors [103]
The conductivity (sdc) of the unblended polyaniline (PAni), PAni–poly(methylmethacrylate) (PMMA) blends and the PAni extracted from blend was measured from room temperature down to millikelvin temperatures at various fields. For PAni (33%)–PMMA (67%), PAni (40%)–PMMA (60%) blends, the reduced activation energy, W ¼ dln(s)=dln(T), decreases on decreasing the temperature below 1 K and the systems are found to be on the metallic side of the MI transition [25a]. In the case of PAni extracted from the blend, the slope change of W occurs at 70 K. For unblended PAni, W increases as temperature decreases (Figure 1.43). The temperature dependence of the conductivity is given by s ¼ s(0) þ Ds1 þ DsL ¼ s(0) þ mT 1=2 þ BT P=2
(1:5)
where the second term arises from the electron–electron interaction and the third term is the correction to the zero temperature conductivity due to localization effect. It is theoretically derived that in the case of electron–photon scattering, p ¼ 3; for inelastic electron–electron scattering, p ¼ 2 and p ¼ 1.5 for the clean and dirty limits, respectively; and very near to the MI transition p ¼ 1. The coefficient of the second term in the above equation is given by 4 3Fs m¼a g 3 2
(1:6)
2 e 1:3 kB 1=2 a¼ 4p2 h 2 hD
(1:7)
where
where D is the diffusion constant and gFs is the interaction parameter. In the case of PAni (40%)–PMMA (60%) blend, there is a deviation from T1=2 dependence above 3 K, whereas in PAni (33%)–PMMA (67%) blend, the T1=2 dependence exists till 4.2 K, which is consistent with metallic behavior near MI transition.
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Polyaniline (as delivered, green)
Absorbance
1
0.75
0.5
0.25
0 200 300 400 500 600 700 800 900 1000 1100
Wavelength [nm] H N+
A−
H A− H−
N
n
Emeraldine base (blue)
Leucoemeraldine base (colourless)
1
1
0.75
0.75
Absorbance
Absorbance
N
H
H
0.5
0.5
0.25
0.25
0 200 300 400 500 600 700 800 900 1000 1100
0 200 300 400 500 600 700 800 900 1000 1100
Wavelength [nm]
Wavelength [nm] N H
N H
N
N
N H
n
N H
N H
N H n
FIGURE 1.42 UV-Vis spectra of polyaniline in different states. (6.16.3) (Reprinted from Wessling, B., Synth. Met., 93, 143, 1998. With permission. Copyright 1998 Elsevier Science)
In the presence of sufficiently high magnetic fields, gmBH >> kBT, s(H, T) ¼ s(H, 0)þm(H) T1=2 where m(H) ¼ a
4 Fs g 3 2
(1:8)
where a, g and Fs are not dependent on the magnetic field [25a]. Below 4.2 K, the conductivity of blends follows a T1=2 dependence and the applied magnetic field reduces the low temperature conductivity (Figure 1.44). The contribution to the magnetoconductivity due to electron–electron (e–e) interactions is given by Ds(H,T ) ¼ agFs T 1=2 0:77 a
gmB kB
1=2
gFs H 1=2
(1:9)
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W = (dln s / dln(T))
10
1
PAni (40%)–PMMA (60%) blend PAni (33%)–PMMA (67%) blend Pristine PAni Extracted PAni from PAni–PMMA blends 0.1 1
T (K)
10
FIGURE 1.43 Log–Log plot of W(T) versus temperature for unblended PAni, PAni (33%)–PMMA (67%), and PAni (40%)–PMMA (60%) blends and extracted PAni. (Reprinted from Rangarajan, G., Srinivasan, D., Angappane, S., and Wessling, B., Synth. Met., 119, 487, 2001. With permission. Copyright 2001 Elsevier Science)
PAni (33%)–PMMA (67%) - 0 T
2.0
PAni (33%)–PMMA (67%) - 5.5 T PAni (40%)–PMMA (60%) - 0 T PAni (40%)–PMMA (60%) - 5.5 T
1.0
s (S/cm)
1.5
fit
0.5
0.0 1.0
1.5
2.0 T 1/2 (K1/2)
2.5
FIGURE 1.44 Plot of s versus T1=2 for PAni (33%)–PMMA (67%) and PAni (40%)–PMMA (60%) blend at zero field and 5.5 T. (Reprinted from Rangarajan, G., Srinivasan, D., Angappane, S., and Wessling, B., Synth. Met., 119, 487, 2001. With permission. Copyright 2001 Elsevier Science)
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0
10
H 2 (T)2
20
30
0.00
1.8 K 3.0 K
−0.01
6.0 K
s (H ) - s (0) (S/Cm)
fit −0.02
−0.03 1.9 K 0.00
3.0 K 6.0 K
−0.01
fit −0.02
−0.03
1.2
1.6 H1/2 (T)1/2
2.0
2.4
FIGURE 1.45 Plot of Ds versus H2 and H1=2 for PAni (33%)–PMMA (67%) blend. (Reprinted from Rangarajan, G., Srinivasan, D., Angappane, S., and Wessling, B., Synth. Met., 119, 487, 2001 With permission. Copyright 2001 Elsevier Science)
gmB 2 Ds(H,T ) ¼ 0:041 a gFs T 3=2 H 2 kB
(1:10)
for high and low magnetic fields, respectively. The magnetoconductivity is found to be proportional to H2 for low fields, gmBH << kBT, and is proportional to H1=2 for high magnetic fields, gmBH >> kBT (Figure 1.45). This proves that the e–e interaction contributes to the magnetoconductivity, which is in the metallic regime. Conclusive additional evidence for the metallic nature of PAni and its blends with PMMA is provided by electron spin resonance (ESR) studies with the observation of a Dysonian line shape [104]. In both cases, the asymmetry ratio (A=B) decreases with decreasing temperature. The observed changes in the line shape from Dysonian to Lorentzian are thus seen to be a manifestation of the variation with temperature of the electrical conductivity (Figure 1.46). The g value is calculated as 2.00191 + 0.00005. The g value, which is close to the free spin value, confirms that the spins are indeed polarons. Resistance versus temperature characteristics for the blends show a trend similar to that of a germanium thermometer (GRT) in the temperature range between 0.35 K and 10 K. The data were fitted using Chebychev polynomials [105]. Sensitivity of the blends is high (better than 0.1 mK) at temperatures below about 1 K, because of the higher resistance. In the temperature range between 2 K and 50 K, the sensitivity is about 1.0 mK. Monotonic trend in MR, where MR ¼
DR R(H,T ) R(0,T ) ¼ R(0,T ) R(0,T )
is similar to the germanium resistor, but the value of MR is less in the case of the blends.
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3.0 A 2.5
20
B 338 D
A/B
H (Gauss) 2.0 10
1.5
s (S/cm)
337 D
1.0 Unblended PAni PAni (33%)–PMMA (67%)
0
0.5 0
50
100
150
200
250
300
T (K)
FIGURE 1.46 Plot of A=B and sdc versus T for unblended PAni and PAni (33%)–PMMA (67%) blend. Inset shows the Diasonian line shape for unblended PAni at 150 K. (Reprinted from Rangarajan, G., Srinivasan, D., Angappane, S., and Wessling, B., Synth. Met., 119, 487, 2001 With permission. Copyright 2001 Elsevier Science)
1.3.2.4.2
Reproducible Fabrication of an Array of Gas-Sensitive Chemoresistors with Commercially Available Polyaniline [106]
Conducting polymers have shown very promising results for application in gas sensors [107] and are currently used in electronic nose systems [108]. Because of the insolubility of these materials, chemoresistors are, in general, prepared by electrodeposition [107a]. Especially considering the limited reproducibility of this method, there remains a need for alternative methods for the preparation of chemoresistors [108b]. The authors of the paper presented a simple, wafer-scale fabrication method based on a commercial polyaniline product (experimental details explained in Ref. [106]). The response behavior of the sensor could effectively be influenced by the choice of the posttreatment process. The presence of noncontinuous rising responses for some sensor types shows that there is not one single response mechanism. This is consistent with the results obtained with more conventional, electrodeposited chemoresistors [107b,109]. Possible interactions are interparticle interaction, swelling [107b,107c], and the modification of the particle resistance because of the partial charge transfer [110]. The baseline resistance of the sensors also depends on the posttreatment. There are various explanations for the mechanism behind the response behavior modification due to the posttreatment. The lower baseline resistance and the faster response transients after the acetone treatment could be caused by dissolving some of the acrylic matrix during this posttreatment. Dissolving the matrix can result in a closer contact between the conducting polymer grains and enhanced diffusion due to a smaller path length through the composite layer. Methanol and ethanol treatments do drastically increase the baseline resistance, which could be caused by the swelling of the polymer inhibiting the charge transfer between the conducting particles. Swelling can be permanent due to only partially reversible sorption of the treatment solvent at room temperature (therefore, the treatment also inhibits the response to the corresponding vapor). The remaining methanol or ethanol could, due to the polar character of these solvents, also be an explanation for the enhanced water response. To illustrate the variance in sensor response behavior, the array was used for vapor recognition, as this is common in electronic noses. The responses of an array of eight sensors, two from each type, were used for a principle component analysis (PCA) of the vapor response. A good separation between the five test
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components was obtained. Even after 2 months of storage of the sensor array, it was, without recalibration, still reasonably capable of distinguishing the five diluted vapors. The methanol and ethanol response patterns suffer from drift, whereas the water, acetone, and isopropanol patterns remain very stable. A simple technique for the batch preparation of gas-sensitive chemoresistors was presented. The stability of the normalized sensor responses was typically 0.2%. An array of such sensors was successfully evaluated for the recognition of a series of diluted organic vapors over a period of 8 weeks. 1.3.2.5 OLEDs Intensive work is currently beingdone to investigate the use of polyaniline as hole injection layers (HILs) in OLEDs, deposited from dispersion.
1.3.2.5.1
PAni as Hole Injection Layer for OLEDs and PLEDs [111]
Since 1990, a large academic and industrial interest in the field of organic light emitting polymers (LEPs) arose from the discovery by Friend et al. [112] that conjugated polymers, the so-called poly[phenylene vinylene]s (PPVs), showed electroluminescence. The new organic diodes consisted of an anode (ITO), a thin PPV layer, and cathodes like Ca, Mg, or Al. With these devices, problems like shortcuts (because of the surface roughness of the ITO) and degradation of the LEP (because of the partial oxidation at the contact area between ITO and LEP) occurred and led to diodes with poor lifetimes. Moreover, the energy gap between the work function of the ITO anode (typically about 4.6–4.9 eV [113]) and the HOMO of the LEP (according to literature, 5.0–5.3 eV [113c,114]) reduced the hole injection from the anode into the polymer. This resulted in poor luminescence values and high driving voltages to obtain light emission. By coating the ITO anode with a HIL, the anode surface could be smoothed and work function could be matched so that the injection barrier for the holes was decreased [115]. So, the use of a HIL in PLEDs has greatly improved the reliability of device manufacturing as well as performance data like light efficiency, driving voltage, and lifetime, but further development is still necessary. Polyaniline (PAni) has been studied for the use as HIL in OLEDs and PLEDs from the early 1990s until today [115a,b,116]. However, there was neither broader scientific research, nor industrial development focused on this topic. Therefore, most of the worldwide display research and development has been carried out with poly(ethylenedioxy-thiophene) (PEDT) as HIL. This is because, for PAni, the most popular approach according to literature was the preparation of solutions [115a,117]. The problems connected with these systems such as bad reproducibility of properties, no long-term stability, and fast gelation may have been the reason that this conductive polymer did not become the focus of industrial development groups. PAni had not been considered as a reliable HIL in commercial devices, until Ormecon’s dispersion technology was transferred into this field. This technology is based on the theoretically well-supported approach of synthesizing an insoluble PAni powder that consists of about 10 nm small primary particles, which can be dispersed in various media like matrix polymers, organic solvents, or water [118,119]. On the basis of this, the joint work between Ormecon GmbH and Covion Organic Semiconductors GmbH that is reported here is the first systematic approach to develop a waterbased, stable, nanoscaled PAni system, which can be easily applied as HIL by spin coating or ink jetting. We will compare the important properties of PAni-based systems to the commercially available PEDT. We present the particle size distribution and resulting film morphology, conductivity, and work function of polyaniline- and PEDT-based dispersions, respectively, HI films as well as the resulting device performance data such as efficiency and turn-on voltage using polymers synthesized by Covion Organic Semiconductors GmbH, as light emitting layers [120]. Water-dispersable, nanoscaled PAni powder was synthesized and dispersed at Ormecon. The particle size distribution was measured by laser Doppler technique with a Microtrac UPA 150 particle analyzer. Investigations on the morphology of PAni films were made at the University of Copenhagen using atomic force microscopy (AFM) and electrostatic force microscopy (EFM). Pictures were made by Tue Hassenkam with a DI Nanoscope IIIa equipped with an extender module.
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The AFM probes were conducting silicon tapping-mode probes with a resonant frequency close to 300 kHz. Conductivity investigations were performed by Fourier transformation electrochemical impedance spectroscopy (FT-EIS) at Ormecon GmbH with an EIS model 6416B. The sheet resistance of PAni and PEDT samples was measured with a two-point probe method using ITO structured glass substrates that were spin-coated with the conductive polymers. Film thicknesses were in the order of 80–100 nm. The measurements were carried out between two coated ITO islands using gold pins as contacts. Impedance spectra were obtained at open circuit potential (E 0 V) by evaluation of perturbation and response signal by FT. Data analysis was carried out using an Rs and C in parallel equivalent circuit (Rs, sheet resistance; C, capacitance of the cell). Further, conductivity measurements were carried out at Covion Organic Semiconductors on specially designed substrates with interdigital structures. Film thicknesses were in the order of 100 nm. The resistance was measured under vacuum to avoid exposure of the somewhat hygroscopic films to humidity. The current was measured for voltages of 10–þ10 V, the resistivity was then calculated from the slope of the straight lines. The reproducibility was very good. The work functions of ITO, PEDT, and PAni layers were determined using a scanning Kelvin probe (SKP, UBM Messtechnik GmbH) in a chamber equipped with silica gel giving a relative humidity of 0%. A Cr or Ni wire with a tip diameter of 80 mm was used as a vibrating reference electrode. The tip was positioned about 20 mm above the specimen, the vibration amplitude was +10 mm, and the vibration frequency of the needle was 1.75 kHz. As measurements could not be performed in ultrahigh vacuum, gold was used as reliable reference material. The work function (which means the HOMO) of the light-emitting polymer, Super Yellow, was measured with cyclic voltammetry (CV). Therefore, the polymer was dissolved in an organic solvent like toluene adding a conductive salt (0.01 mol=l TBATFB) to the solution. Using a Pt working and counter electrode and an Ag or AgCl reference electrode, the oxidation potential of the LEP could be determined. The HOMO energy for Super Yellow was calculated from the oxidation potential. To check the PLED device performance, PLEDs in the standard configuration ITO=20 nm PAni=80 nm LEP=6 nm Ba=100 nm Ag were prepared and measured by increasing the voltage while detecting the current and the luminance. The commercially available, stable polyaniline dispersion for the spin coating of HILs—the developmental type Covion PAT 010—has a particle size distribution of the water-based PAT 010 compared to commercially available PEDT samples as shown in Figure 1.47. It is evident that the mean particle size of the PAni dispersion is at least a factor of 6 smaller than for PEDT. Although the number distribution for the nanoscaled PAni dispersion shows a maximum at about 35 nm in particle size, this is the average particle size for the PEDT sample with the smaller particles around 200 nm. An additional advantage of PAT 010 is that the PAni content in this nanosystem can be varied between 1 and 7 wt % without affecting the particle size distribution. The small size distribution of PAni is stable and stays the same at least over a 12 months test period. Spincoating the PAni dispersion onto ITO or glass, very smooth HI films without significant surface roughness can be obtained. Morphology results are shown in the following AFM and EFM pictures (see Figure 1.48). This very homogenous surface of the dried PAni layer with a maximum roughness <10 nm proves that the PAT 010 dispersion is well suitable for the fabrication of smooth HILs and guarantees optimum wetting and adhesion contact to the following LEP layer. Even thinner PAni films of about 60 nm thickness are sufficient to flatten the surface roughness of ITO to avoid direct contact of ITO with the LEP layer, greatly improving the reliability of device manufacturing. The morphology combined with a low lateral conductivity of the PAni films allows the production of crosstalk-free passive matrix displays with improved performance entirely made by spin-coating. Depending on the acid content in the water-based PAni dispersions, the conductivity of the HILs can be varied between 106 and 102 S=cm. In order to avoid crosstalk in passive matrix applications, the conductivity of PAT 010 has to be below 105 S=cm. The determined conductivity values depend
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24
24 PAni PAT 010.au - 30 nm PEDT sample 1.au - 200 nm PEDT sample 2.3u - 400 nm
22
Absorption [%]
20
22 20
18
18
16
16
14
14
12
12
10
10
8
8
6
6
4
4
2
2
0
0 1E
0.01
−3
0.1
1
10
Particle size [µm]
FIGURE 1.47 Comparison of the particle size distribution of a water-based standard PAni dispersion (PAT 010, patterned columns) with standard PEDT samples (gray and dark gray columns) measured by laser Doppler technique with a Microtrac UPA 150 Particle Analyzer. (Reprinted from Werner, B., Posdorfer, J., Wessling, B., Becker, H., Heun, S., Vestweber, H., and Hassenkam, T., Proc SPIE 4800, 115, 2003. With permission. Copyright 2003 SPIE)
strongly on the measuring method used, as can be seen in the following investigations. Measurements at Ormecon were performed on glass substrates having four ITO islands. The substrates were coated with PAni and PEDT samples and then conditioned for 20 h at a well-defined humidity. The conductivity of the PAni and PEDT layers was determined as described in the experimental part using FT electrochemical impedance spectroscopy. A comparison of the humidity dependent conductivity of PAT 010 and commercially available PEDT dispersions is shown in Figure 1.49.
10.0
10.0
0.1 V
10.0 nm 7.5
7.5
5.0
5.0
5.0 nm
2.5
2.5
0.0 V
0.0 nm 0
2.5
5.0
7.5
0 10.0 µm
0.1 V
0
2.5
5.0
7.5
0 10.0 µm
FIGURE 1.48 AFM (left) and EFM (right) pictures of PAT 010: a 100 nm thick PAni based HIL spin coated onto glass and dried at 1808C Pictures made by Tue Hassenkam at the University of Copenhagen. (Reprinted from Werner, B., Posdorfer, J., Wessling, B., Becker, H., Heun, S., Vestweber, H., and Hassenkam, T., Proc SPIE 4800, 115, 2003. With permission. Copyright 2003 SPIE)
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−0.5 *
−1.0
* *
*
*
Log conductivity
−1.5 −2.0 −2.5 −3.0 −3.5 −4.0 *
PAni PAT 010 PEDT sample 1 PEDT sample 2
−4.5 0
10
20
30
40
50
60
70
80
90
100
Relative humidity [%]
FIGURE 1.49 Conductivity in dependence on relative humidity measured by FT EIS. ITO structured glass was spin coated with PAni and PEDT and dried at 1808C leading to coating thicknesses of 80–100 nm. (Reprinted from Werner, B., Posdorfer, J., Wessling, B., Becker, H., Heun, S., Vestweber, H., and Hassenkam, T., Proc SPIE 4800, 115, 2003 With permission. Copyright 2003 SPIE)
Although PAni and PEDT sample 1 show a strong dependence of the conductivity on the relative humidity, the conductivity of PEDT sample 2 remains high and almost constant over the entire humidity range. Moreover, there is an obvious difference between the conductivities at 0% humidity. After a conditioning time of 20 h, PAni has a low conductivity of 104 S=cm, whereas PEDT sample 1 shows a medium conductivity of 103 S=cm. Contrary to those two low conductive layers, PEDT sample 2 is by far more conductive with values of about 0.3 S=cm. This means that this particular PEDT is three orders of magnitude more conductive than PAni. Long-term investigations of the conductivity development at constant humidity have shown that the PAni-based HILs are stabilizing their sheet resistance toward lower values during a storage time of >20 h. Figure 1.50 shows the film conductivity over a test period of 10 d. For the constant humidity of 90%, the conductivity is decreasing linearly within the first 24 h. After one day of conditioning, the film has almost reached its equilibrium with a final conductivity of 5 103 to 102 S=cm. Equilibration time for the sample under 50% humidity takes longer. This curve shows an asymptotic behavior, reaching a constant value of 6 104 S=cm after 5–6 d. The values for the conductivity at 0% humidity are extrapolated (using the values for 90%, 70%, 50%, and 35% RH) and show as well an asymptotic behavior. This PAni system is equilibrated after a 7 d storage time ending with a conductivity of 2–3 105 S=cm at 0% RH. Additional PAni-based systems as well as PEDT sample 1 are under investigation. Lateral conductivities of several PAT 010 batches were also measured at Covion. The samples were spin-coated onto an interdigital ITO structure, allowing to measure the resistivity with a voltmeter (2 point probe measurement). Measurements (under vacuum) for different batches resulted in values between 3.2 106 and 6.5 106 S=cm. The presented investigations considering the conductivity of PAni and PEDT films have been carried out using different measuring methods as well as humidity conditions. We can conclude that PAni and PEDT
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Log Lf
Conductive Polymers as Organic Nanometals
−1.0
−1.0
−1.5
−1.5
−2.0
−2.0
−2.5
−2.5
−3.0
−3.0
−3.5
−3.5 −4.0
90 % Relative humidity 50 % Relative humidity
−4.5
0 % Relative humidity −5.0
−5.0 0
50
100
150
200
250
Storage time [h]
FIGURE 1.50 Conductivity development of PAni-based HILs in dependence on storing time under constant humidity conditions and ambient pressure. The PAni sample was spin coated onto ITO structured glass and dried at 1808C, coating thickness 100 nm. (Reprinted from Werner, B., Posdorfer, J., Wessling, B., Becker, H., Heun, S., Vestweber, H., and Hassenkam, T., Proc SPIE 4800, 115, 2003 With permission. Copyright 2003 SPIE)
sample 1 principally show the same dependence of the conductivity on the chosen humidity. The best method to accelerate the equilibration of the HIL at the OLED processing humidity of 0% RH is to measure the conductivity of the thin films under vacuum. Due to the similarity with the geometry of the final application on a passive matrix substrate, this method is the best way to measure the film conductivity of HILs for passive matrix displays where the lateral conductivity has to be low to avoid crosstalk problems. Device performance is influenced by morphology aspects as well as by hole injection into the LEP film where the recombination of holes and electrons causes light emission. To make efficient PLEDs with high luminescence values at comparatively low operating voltage (3–5 V) it is necessary to optimize charge injection from the anode into the LEP layer. Using PAni or PEDT as HIL, the injection barrier between the ITO anode and the polymer film can be decreased and luminescence can be improved. The following investigations of the work function of pure ITO and PAni or PEDT coated ITO were carried out using the scanning Kelvin probe method. As light emitting layer, the commercially available Super Yellow (from Covion Organic Semiconductors [120]) was used. The HOMO energy of Super Yellow was determined using cyclic voltammetry. Figure 1.51 shows a comparison of values found for Super Yellow and the different anodes. The percentage values shown on the y-axis of PAni and PEDT sample 1 coated ITO are calculated by referencing the maximum EL efficiency of a Super Yellow device against the standard HIL material PEDT sample 2 (PEDT 2 ¼ 100%). Comparing the work functions of PAni and PEDT with the corresponding performance values, it appears that there are three different groups of PAni dispersions: <100%, 100%, and >100% performance compared to PEDT sample 2. PAni films made with the commercially available PAT 010 are at least equal to PEDT sample 2. Steadily increasing the dispersion
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Optimum range
Performance [%] Super yellow
PAni < 100%
PAT 010. 100%
PAni > 100%
PEDT #197% PEDT #2 = 100% OC PPV 4.2
4.3
4.4
4.5
4.6
4.7
4.8
4.9
5.0
5.1
5.2
5.3
5.4
Work function [eV]
FIGURE 1.51 Work function determined with the scanning Kelvin probe method (used for ITO and PAni or PEDT coated ITO, coating thickness 80–100 nm) and cyclic voltammetry (used for LEP Super Yellow synthesized by Covion Organic Semiconductors [120]). Corresponding performance values in percent were obtained from the Super Yellow maximum EL efficiency of a comparable PEDT device made in the same layer configuration (PEDT sample 2 ¼ 100%). (Reprinted from Werner, B., Posdorfer, J., Wessling, B., Becker, H., Heun, S., Vestweber, H., and Hassenkam, T., Proc SPIE 4800, 115, 2003 With permission. Copyright 2003 SPIE)
quality and the work function into the optimum range of 4.85–4.95 eV could stabilize device performance. With PAT 010, device performance now lies reproducibly at least at 100%, which means the waterbased PAni dispersions are very well suitable as HILs. PLED performance was measured on test devices manufactured at Covion, using either Covion’s PDY 132 (commonly called Super Yellow) or Covion’s polyspiro materials [120] for other colors as the active polymer layer on 20 nm PAT 010 PEDT as HIL. The efficiency measurements clearly showed that PAni PAT 010 improves the device performance for Polyspiro Blue (by 13%) and Deep Green (by 26%). Using Super Yellow as LEP, PAT 010 is equal to the standard PEDT (sample 2). Regarding the stability of the PAni dispersions we have already mentioned that the particle size distribution is stable for more than one year. The following investigation (Figure 1.52) shows the dispersion stability in terms of device performance for freshly made Super Yellow devices with 20 and 100 nm PAT 010 as HIL. The maximum efficiency for the 20 nm thick PAT 010 remains almost constant at 100% over the whole tests period of 12 months. It is remarkable that the luminescence for the 100 nm thick PAni based HIL increases by about 20% within the first 2 months, remaining constant from there on. This means that PAT 010 renders consistently good devices over a storage period of more than 1 y and so shows excellent long-term stability when stored at room temperature. The synthesis of a nanoscaled, water-dispersible polyaniline powder leads to long-term stable, agglomerate-free HI system that can be easily spin-coated onto ITO. The dried PAni layers show significantly lower conductivity than PEDT (105 with EIS respectively 5 106 S=cm on
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Performance in %
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140
140
120
120
100
100 Start
1y
80
80
60
60
40
40 PAT 010, 20 nm PAT 010, 100 nm
20
20
0
0 0
2
4
6
8
10
12
Storage time for PAT 010 [months]
FIGURE 1.52 Performance stability in dependence on storage time of a PAT 010 dispersion. PLEDs manufactured at Covion using PAT 010 (20 and 100 nm) as HIL and Super Yellow as LEP (80 nm). Performance values in percent are calculated using the maximum EL efficiency value for 20 ref. 100 nm thick PAT 010 at storage time ¼ 0 mon. (Reprinted from Werner, B., Posdorfer, J., Wessling, B., Becker, H., Heun, S., Vestweber, H., and Hassenkam, T., Proc SPIE 4800, 115, 2003 With permission. Copyright 2003 SPIE)
interdigital structures) and prevents the occurrence of crosstalk in passive matrix displays. The most important improvement compared to commercially used PEDT is for sure the decrease of the injection barrier between ITO anode and polymeric emitter. By increasing the work function to be close to the energy level of the LEP, the device performance could be stabilized at 100% (compared to PEDT standard HI material) for Super Yellow and significantly improved for the blue (13%) and green (25%) light-emitting polymers when using PAni as HIL. The stability of the dispersion properties like stable particle size distribution and constant luminescence data for the spin-coated and dried HIL is excellent, leading to a reliable PAni material. As PAT 010 is one of Covion’s commercial products under steady development, there will be further improvements in efficiency and operating voltage in the nearby future. First steps concerning the production scale-up from laboratory scale to technical scale without affecting the dispersion and HIL properties have already been successfully realized. Investigations on device lifetime are running. First results have shown that lifetime is not reduced by using PAni instead of PEDT. This means that PAT 010 is a reliable alternative to the widespread used PEDT.
1.3.2.5.2
Characterization of Polyaniline-Based Polymer Light-Emitting Devices during Operation by Electrical Impedance Spectroscopy [121]
Polyaniline was used as hole injection material for polymer light-emitting devices (PLED) spin-coated from a water-based polyaniline=poly-styrenesulfonate (PAni=PSS) dispersion. Data presented here were based on red, yellow, green, and blue LEPs. Depending on the applied bias voltages, the devices were studied by electrical impedance spectroscopy (EIS) in a wide frequency range. Experimental data could
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only be fitted by applying an equivalent circuit consisting of three RC elements in series. By separating the frequency-dependent bulk and interface contributions, the bulk and junction resistance and capacitance were determined for various current densities. Variation in high interfacial capacitance was observed in the devices before and after stressing the PLEDs with a constant current of 10 mA=cm2 for 100 h. Polymeric light emitting diodes have emerged as one of the most promising technologies for inclusion in flat-panel displays. They have various advantages for high display quality such as lightweight, wide viewing angle, bright emission, and rapid response. However, PLED device performance is still strongly required to improve power efficiency, long lifetime, and high color purity. Unfortunately, most lightemitting polymers still lack a satisfactory stability, especially the blue. Impedance spectroscopic investigations on PPV-based LEDs are reported in the literature by several groups with different results and interpretations. Impedance spectra on ITO=PPV=Al devices were described by two semicircles within a Schottky model representing bulk and junction region [122] or by the presence of an interfacial oxide layer at the PPV=Al contact [123]. Equivalent circuits using three RC elements suggested a spatial variation of the conductivity in the PPV film [124]. A more complex equivalent circuit was proposed by analyzing the Poisson’s and the hole and electron continuity equations [125]. Here we present a detailed study of PLEDs with different LEPs by electrical impedance spectroscopy in a wide frequency range as a function of current corresponding to the applied bias voltage. From the frequency dependence information on the transport parameters of the red, yellow, green, and blue lightemitting materials and the equivalent circuit necessary to describe the device characteristics are expected. The current dependence should give information about changes in interfacial and bulk regions. The light-emitting devices were prepared by Covion by spin coating and curing a 80 nm layer of PAni=PSS as HIL onto indium tin oxide (ITO) patterned glass substrates followed by spin coating of the 80 nm LEP layer. A water-based PAni=PSS dispersion (Ormecon PAT020) with a conductance of 1 105 S=cm and a work function of 5.1 eV was used [126]. Light-emitting polymers were red (AEF 2145), yellow (PDY 132), green (AEF 2394), and blue (AEF 6053) from Covion Semiconductors GmbH [127]. As cathode material, 6 nm Ba with a capping layer of 100 nm Al for lowering cathode resistance was deposited by evaporation. Finally, the device is protected from water and oxygen using a glass lid. The average electrode area was 16 mm2. For electrical impedance spectroscopy and measurement of I–V characteristics, an Autolab potentiostat=galvanostat model PGSTAT 30 was used. For EIS, a digital signal converter, a signal conditioning unit, and a fast analog to digital converter with two channels (Autolab, model FRA 2) were used in a frequency sweep range from 10 mHz to 1 MHz. An AC perturbation signal of 50 mV was applied upon constant DC forward-bias in the range of 0–5 V. During the measurements, the devices were housed in a Faradaic cage. Impedance data were fitted to an equivalent circuit using simulation and modelling software based on the work of Boukamp [128]. All experiments were performed at room temperature. The PLED devices were stressed at 10 mA=cm2 applied by a Keithley source meter, model 2400. In order to evaluate experimental data, an equivalent circuit consisting of three parallel RCs in a serial combination with a contact resistance Rc, as shown in the inset of Figure 1.53, was used. In the equivalent circuit, the interfacial layers correspond to the junction resistance R1 and R2 and capacitance C1 and C2. Between these depletion layers exists a bulk region represented by a resistor Rb and a capacitance Cb. For the devices, a model was assumed where the PAni=LEP and the LEP=Ba interfaces were regarded as Schottky barriers. The contact of LEP with the high work function PAni was regarded as ohmic but the interaction of LEP with the conducting polymer can result in the creation of defect states forming an interfacial region between hole injection layer and the LEP [129]. Similar effects have been reported for hole injection from an Ag anode into a dialkoxy-PPV explained by electron trapping near the anode [130]. At the cathode doping [131], surface contamination [132] and different Fermi levels can be responsible for the formation of a Schottky contact at the LEP–metal interface [72]. Band bending at the interfaces are accompanied by an inhomogeneous electric field distribution and depletion layers occur at
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4000 Rc 3000
R1
Rb
R2
C1
Cb
C2
40 nA 50 nA
−Zim [kΩ]
95 nA 2000
194 nA 488 nA
1000
0 0
1000
2000
3000
4000
Zre [kΩ]
FIGURE 1.53 Impedance Cole–Cole plots of unstressed Pani=Green=Ba device at various currents showing experimental (symbols) and fitted (solid line) data. The left semicircles were assigned to the polymer bulk and the right semicircles to the depletion layers [121].
the interfaces. Between the Schottky barriers a space–charge region in the bulk is formed. It is also important that both electrodes inject carriers that can dominate the transport. By fitting of the impedance data, the resistance Rb and the capacitance of the bulk material Cb and the interfacial resistance R1 and R2 and capacitance C1 and C2 were obtained. A contact resistance of about 30 V independent of applied bias voltage was approximately the same for all devices. The dependence of the capacitance on the current corresponding to the applied bias voltages is shown in Figure 1.54. The capacities of the interfacial regions were significantly higher than the capacitance in the bulk. The sum of the resistance values Rc, Rb, R1, and R2 agreed very well with the numerical derivatives of the corresponding I–V curves of the devices. We believe that C2 can be related to the depletion area at the PAni or LEP interface. The magnitude of the capacitance C2 was much higher than the bulk capacitance and the capacitance C1 that was related to the LEP–Ba interface. This was proven by variation of hole injection layer thickness. For a given device, the values for C2 changed proportional to hole injection layer thickness. When the cathode consisted of LiF=Al only changes in C1 were obvious. As shown in Figure 1.54, for a yellow device, the capacitance of Cb was almost independent of current density. For C1 and C2, an increase of capacity was observed for the unstressed and stressed devices. At higher current densities, the capacitance decreased to the initial values forming a current range for high capacitance. For different LEPs, these ranges are summarized in Table 1.3. For red, only a slight variation in capacitance was observed. For all other devices this broadening increased in the sequence yellow, green, and blue. Impedance data for PLEDs could only be fitted by applying an equivalent circuit consisting of three RC elements in series. By separating bulk and interface contributions, bulk and interfacial capacitances were identified and determined at various current densities. Current ranges with high capacitance were observed in devices based on yellow, green, and blue LEPs. The highly capacitive junction provides interfacial electron–hole recombination, which can lower the lifetime and efficiency of the displays. It is important that the charge mobility is balanced in the devices to maximize power conversion efficiency
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10000 0h 1000
Capacitance [nF]
100 10 1 0.001
0.01
0.1
1
10
100
1000
10000
10000 100 h 1000 100 10 1 0.001
0.01
0.1
1
10
100
1000
10000
Current [µA] C1
Cb
C2
FIGURE 1.54 Capacitance for bulk and interfacial regions for stressed and unstressed PAni=Yellow=Ba device as a function of current [121].
and minimize degradation of the material by oxidation. As the range for high capacitance was broadest for the blue light-emitting material, the unstable electroluminescence of this material can be better understood. 1.3.2.6 Polymer Electronics (Research: OLEDs, FETs) Ormecon GmbH offers a wide range of dispersions of the organic metal polyaniline for use in basic and applied research for polymer electronics in universities and other public research institutions.10 A series of dispersions is available based on water or a variety of organic solvents, covering a wide range of conductivity between 102 and 107 S=cm. These are based on Ormecon’s unique and proprietary synthesis and dispersion technology. An introductory overview covering the scientific basis and further references can be found in Ref. [133]. An overview describing the patent portfolio is available as well.11 TABLE 1.3 Current Range for High Capacitance in Interfacial Regions for Different Light-Emitting Polymers before and after Stressing at 10 mA=cm2 Current Range [mm) for LEP Operating Time [h] 0 100
Red
Yellow
Green
Blue
10–100 10–500
1–100 0.05–1000
0.04–100 0.01–1000
1– 10000 0.01– 10000
10 There are other products available—dispersions, paints, coatings, blends—for industrial commercial applications or R&D use outside of polymer electronics: printed circuit board final solderable finish; antistatic coatings; screen printable electrodes; antistatic and anticorrosion coatings. 11 Ask for ‘‘Condensed description of the Zipperling or Ormecon patent portfolio: range of covered areas.’’
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Conductive Polymers as Organic Nanometals TABLE 1.4
Technical Data Survey for D 1005 W
D 1005W
Grade A
Conductivity of the layer Work function of the layer Particle size Aniline unit: sulfonic acid group Solid content Max solids possible pH Viscosity Color Spin coating data for both grades D 1005 W
5
10 S=cm 5.0 eV <100 nm 1:3:4 4% (7%) 1 4 mPa s green 5s 30 s drying recommended layer thickness
Grade B 2
10 S=cm 5.0 eV <100 nm 1:2 4% (7%) 1 3 mPa s green 500 rpm 3000 rpm 10 min at 1808C 50 nm (140–200 nm possible)
Ormecon GmbH offers an active support for your research—we will try our very best to change formulations to better fit your requirements or help you realize your ideas. However, we suggest that you start working with a product that is regularly and reproducibly available. All these products are precommercial trial products. These are the following: D1005 W A water-based dispersion of polyaniline recommended for use as HIL in OLEDs (see Table 1.4). Two grades are routinely available: . .
Grade A: conductivity 105 S=cm Grade B: conductivity 102 S=cm
D 1020=D 1024 Xylene based dispersions allowing preparing thin layers or films with 100 and 200 S=cm (see Table 1.5). Depending on the required application, different dilutions are available described by respective product numbers. We highly appreciate your specific inquiry to provide you with advice and technical data for the best suitable dispersion. 1.3.2.7 Electroluminescence (EL) Displays [134] Electroluminescence (EL) lamps are so-called Lambert emitters, i.e., the density of the light emitted from the surface is identical from all sides. The light is very narrowband, almost monochromatic, very homogenous, and in the visible range. These are the ideal parameters for advertising and security applications, and opens new prospects that would not be possible, or would demand much higher investments, with traditional technologies.
TABLE 1.5
Technical Data Survey for D 1020 and D 1024
Dispersion Conductivity of the layer Particle size Solid content Viscosity Spin coating data
D 1020
D 1024
200 S=cm 150 nm medium 2,25% 5,4 cP 5 s 500 rpm 5 s 1500 rpm drying 120 s at 1008C
350 S=cm 15 nm medium 2,1% 8 cP
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Not illuminated
FIGURE 1.55
Illuminated
EL display used in a sign.
ORMECON EL displays are thin, lightweight, and have uniform surface illuminations that can be applied on almost any substrate. They are impact and vibration resistant, flexible, i.e., can be applied on flexible substrates, and require low power consumption because they generate almost no heat in operation. Today EL displays are already used for several applications such as (see Figure 1.55 through Figure 1.58) . . . .
Advertisements Automotive and aircraft applications (instruments, functional and decorative illuminations Portable consumer electronics Architectural applications
The market for EL displays is constantly growing because of the more demanding requirements of the advertisement industry, consumer electronics, and automotive industry. Such requirements are of technical, as well economical nature. Complete ready-to-use EL displays with custom design including all necessary power supply can be manufactured (see Table 1.6). Not only full or part area illuminations but also animated designs have been realized. 1.3.2.8 Dramatic Increase in Conductivity [135] Some applications for conducting polymers (like source, drain, gate, and connector in organic field effect transistors, [transparent] electrode or raw material for highly conducting paints and lacquers with
FIGURE 1.56
Electroluminescence (EL) displays for advertising applications [133].
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Conductive Polymers as Organic Nanometals
FIGURE 1.57
1-67
EL display in application (first pilot project with AVIS—a large German rental car company) [134].
a view on to EMI shielding) demand a very high conductivity (>102 S=cm). To address these applications, we focussed on improving the conductivity of polyaniline dispersions. As the conductivity of raw polyaniline powder is about 5 S=cm, one needs to improve this conductivity by a factor of 20–100 to be able to find an application in the above-mentioned fields. The solution to the problem of increasing the conductivity to above 102 S=cm, and doing this in a reproducible way has not been found in another variation of ‘‘EB=special dopant=secondary dopant= stretching’’, but on the following basis: .
.
Synthesis of a dispersible powder according to Ormecon’s technology [136], exhibiting a conductivity of 2–5 S=cm Use of p-toluene sulfonic acid as the dopant (both in PAni and PEDOT)
FIGURE 1.58
Close up of sign shown in Figure 1.57 [134].
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TABLE 1.6
Technical Information for EL Displays
Thickness Color=Animation Working Temperature Storage Brightness Consistency Wavelength Luminance Luminosity
Down to less than 1 mm On request 20 to þ458C (high temperatures reduce the luminosity and life expectancy) Dry=cool=dark (10 to þ258C) More than 95% Approx. 420–750 nm (according to the pigment) >100 cd=m2 , @135 V and 650 Hz (pink=white) >250 cd=m2 can be achieved at high frequencies (at highest luminosity, lifetime will be affected) Half-time ranges from 1.000 to 10.000 h depending on the light emissive pigment used and depending on luminosity Very high, cold light source (no infra-red, no UV) 110130 V AC (standard) Sinus, frequency ¼ 100–1.000 Hz In tests @ 130 V AC, 400 Hz approx. 35 W=m2
Lifetime Effectiveness Power Supply Frequency Power Consumption
. . .
Predispersion to about 40–50 S=cm [137] Subsequent treatments and additional dispersion [up to 500 S=cm] Final formulation to achieve best film forming and spin coating properties
Hence, the conductivity of the original powder has been increased two times each by a factor of 10, without changing chemistry (as can be seen in the IR spectra, Figure 1.59). It should be noted that not only special dopants, but also secondary dopants are unnecessary, as we achieved such conductivity also with octanole as dispersion medium. This complex sequence of process steps is controlled by at least 80 different parameters (like temperatures, time, sequence of actions, additives, etc.). We have analyzed the effect of all of these parameters from the first to the last process step levels and determined the most important 25 parameters. These were the input into a neural network program [138]; the neural network has been appropriately set up and optimized, and has been trained in several series of training actions finally with the input of more than 400 fully documented experiments. A simplified flow scheme of the final results describing the effects of the various parameters is shown in Figure 1.60. It can be seen that some parameters have a linear (positive or negative) influence on the
57
52
47 5S-PAni 180S-PAni 42 4350
3350
2350 Wavenumber [cm–1]
FIGURE 1.59
IR Spectrum [135].
1350
350
Transmission [%]
62
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Conductive Polymers as Organic Nanometals
1B
1O
1E
1C
Legend:
1D
Positive influence Negative influence
1T
Level 1 (1W) 1U
1R
Nonlinear
2C
2D
Local Minimum
Level 2 (2R) 2F
2E 2G
2H 2K Level 3 (3H) 3C
3B
3F 3D Level 4 (4J) 4E
4D
5D2 Level 5 (5G)
5D1
5B 5C1 5C2
FIGURE 1.60
5F
5E
Flow diagram with parameter interactions [135].
ultimate output value (conductivity), some have a nonlinear effect, and some show a local minimum. It became evident that the parameters in the whole process are highly interconnected and have nonlinear interactions in a multidimensional manner. It is important to note that the conductivity values reported here are measured on films and layers, which have been spin-coated or deposited from dispersion. These samples have not been stretched. The conductivity is isotropic. The crystallinity, determined by x-ray, does not seem to be increased. The intrinsic (AC) conductivity of the metallic core in the primary particles has been determined to be 103 S=cm (10 GHz and 9.7 GHz, respectively) [30]. It will be the object of the further research whether these values can also be achieved under DC conditions, and whether 103 S=cm is the limit of conductivity for this system. In addition, screen printable pastes are available, which can be used for printing back and front electrodes for electroluminescent films. Table 1.7 shows technical data for the first EMI shielding paint and EL electrode printing paste. Actually, the most interesting and most surprising potential application is for the manufacturing of electrode materials for the so-called supercapacitors. Here, such dispersions have been used for the preparation of electric double layer capacitors. Such capacitors have shown a dramatically increased performance compared to any other known energy storage devices, as can be seen in Figure 1.61. The energy density and the power density is bigger by a factor of 10 to 100 compared to conventional supercapacitors and even Li ion or Ni–Cd batteries. TABLE 1.7
Technical Description of EMI and EL Paints EMI Shield Paint
Solvent base
xylene
Dilution with Solid content Application Layer thickness Specific conductivity Surface resistance
xylene 8.5% spray coating 2–5 mm 150 S=cm <10 V=sq
EL Screen Print Paste aromatic solvent mixture (solvent naphta=xylene based) xylene 15% screen printing 7 mm 70 S=cm 50 V=sq
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Wh/kg
Conjugated Polymers: Processing and Applications
0.5 mA/cm2
1k Battery
1.0 mA/cm2
Evaluated capacitor
5.0 mA/cm2
Initial value Li-ion
100
NiWH Ni-Cd
10
After 4000 cycle
Super Capacitor Condensor 50F
AI
Coin
10
FIGURE 1.61 [135].
100
1k
10k
W/kg
Ragone Diagram: Efficiency of new supercapacitor type (Figure provided by Nissan Chemical Ind.)
A new application, which might have the chance to become realized more in short term is in the area where Ormecon has successfully introduced a new nanotechnology on a worldwide basis: the printed circuit board industry and here, more specifically, solderable surface finishes. The product is a new allorganic solderable surface nanofinish for printed circuit boards (which can be introduced into the markets within a relatively short time). The new products allow producing a solderable surface in a onestep process.
1.3.3
Final Remarks
There will be many more industrial applications in the future, as soon as the new useful property combinations of organic metals, especially polyaniline, are more broadly known, and more confidence has been built in the market based on the first pioneering applications. There might come sensors, actuators, or gas separation membranes, or even totally different applications that we even now cannot imagine. However, several potential applications are actually not seriously approached, even though their principal feasibility has been shown. It seems that industrial development groups refrain from working with PAni as long as they are conceptually and mentally biased towards either a direct polymerization approach or a solution technique. The development of electrochromic windows, sensors, and gas separation membranes would need a reproducible nanotechnology for applying PAni to the substrate as well as progress in electrolytic and double layer capacitors and solar cells. Our dispersion concept and theory could help to speed up such development work, but only if industrial laboratories interested in the development and marketing of new technologies based on conductive polymers (or organic metals) like polyaniline become open to the use of dispersion processed PAni. With two basically different and long-term industrial uses of polyaniline, in corrosion protection and in the manufacture of printed circuit boards, both based on and only feasible with dispersion technologies, the dispersion concept has shown its value and has allowed the first significant commercial application of conductive polymers. A corresponding scientific and theoretical understanding has also been elaborated [15] and is worth being discussed and investigated much more intensively. Organic metal nanotechnology has begun.
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93. (a) Putilova, I.N., S.A. Balezin, and V.P. Barannik. 1969. Metallic corrosion inhibitors. New York: Pergamon; (b) Poling, G.W. 1970. Corros Sci 10:359; (c) Roberts, R.F. 1974. J Electron Spectrosc Relat Phenom 4:273. 94. (a) Wessling, B. 1994. Adv Mater 6:226; (b) Brusic, V., M. Angelopoulos, and T. Graham. 1997. J Electrochem Soc 144:436. (c) Wessling, B. 1999. Circuit World 25:8. 95. (a) Andra¨, K. 1996. Metalloberfla¨che 50:176; (b) Bratin, P., M. Pavlov, and G. Chalyt. 1999. Print Circuit Fabr 22:30; (c) Tench, M., M.W. Kendig, D.P. Anderson, D.D. Hillmann, G.K. Lucey, and T.J. Gher. 1993. Soldering Surf Mount Technol 13:18. 96. Ladebusch, H., T. Strunskus, J. Posdorfer, and B. Wessling. 2001. Synth Met 121:1317. 97. Wessling, B. 1998. Synth Met 93:143. 98. (a) Lu, W.K., R.L. Elsenbaumer, and B. Wessling. 1995. Synth Met 71:2163; (b) Ahmad, N., and A.G. MacDiarmid. 1996. Synth Met 78:103; (c) Fahlmann, M., S. Jasty, and A.J. Epstein. 1997. Synth Met 85:1323; (d) Schauer, T., A. Joos, L. Dulog, and C.D. Eisenbach. 1998. Prog Org Coat 33:20. 99. Strunskus, T., M. Grunze, G. Kochendoerfer, and Ch. Wo¨ll. 1996. Langmuir 12:2712. 100. Lin, S.L., K.L. Tan, and E.T. Kang. 1998. Langmuir 14:5305. 101. Arendt, N., F. Baron, V. Benz, M. Letterer, H. Merkle, S. Schroeder, and B. Wessling. OnBoard Technology pp. 30–32, Wise Media. 102. Hugot–Le Goff, A. 1997. Handbook of organic conductive molecules and polymers, vol. 3, ed. H.S. Nalwa, 745–782. Chichester: Wiley. 103. Rangarajan, G., D. Srinivasan, S. Angappane, and B. Wessling. 2001. Synth Met 119:487. 104. Srinivasan, D., T.S. Natarajan, G. Rangarajan, S.V. Bhat, and B. Wessling. 1999. Solid State Commun 110:503. 105. Bharthwaj, A.N., S. Angappane, D. Srinivasan, T.S. Natarajan, G. Rangarajan, and B. Wessling. 2000. 18th International Cryogenic Engineering Conference 2000 held at IIT Bombay during 21–25 February 2000. 106. Meijerink, M.G.H., D.J. Strike, N.F. de Rooji, and M. Koudelka–Hep. 2000. Sens Actuators B 68:331. 107. (a) Bartlett, P.N., and K. Ling–Chung. 1989. Sens Actuators 20:287; (b) Slater, J, M., E.J. Watt, N.F. Freeman, I.P. May, and D.J. Weir. 1992. Analyst 117:1265. 108. (a) Pearce, T.C., J.W. Gardner, S. Friel, P.N. Bartlett, and N. Blair. 1993. Analyst 118:371; (b) Mielle, P. 1996. Lectronic noses: Towards the objective instrumental characterization of food aroma. Trends Food Sci Technol 7:432. (c) Gardner, J.W., and P.N. Bartlett. 1999. Electronic noses, principles, and applications. Oxford: Oxford Science Publication. 109. Patridge, A.C., P. Harris, and M.K. Andrews. 1996. Analyst 121:1349. 110. (a) Josowicz, M., J. Janata, K. Ashley, and S. Pons. 1987. Anal Chem 59:253; (b) Blackwood, D., M. Josowicz. 1991. J Phys Chem 95:493. 111. Werner, B., J. Posdorfer, B. Wessling, H. Becker, S. Heun, H. Vestweber, and T. Hassenkam. 2003. Proc SPIE 4800:115. 112. Burroughes, J.H., D.D.C. Bradley, A.R. Brown, R.N. Marks, K. Mackey, R.H. Friend, P.L. Burn, and A.B. Holmes. 1990. Nature 347:539. 113. (a) Campell, A.J., D.D.C. Bradley, and H. Antoniadis. 2001. J Appl Phys 89:3343; (b) Kugler, T., and W.R. Salaneck. 1999. Chem Phys Lett 310:319. (c) Cacially, F., R.H. Friend, N. Haylett, R. Daik, W.J. Feast, and D. Santos. 1996. Appl Phys Lett 69:3794. 114. Campbell, I.H., T.W. Hagler, D.L. Smith, and J.P. Perraris. 1996. Phys Rev Lett 76:1900. 115. (a) Yang, Y., and A.J. Heeger. 1994. Appl Phys Lett 64:1245; (b) Heeger, A.J., I.D. Parker, and Y. Yang. 1994. Synth Met 67:23. (c) Carter, S.A., M. Angelopoulos, S. Karg, P.J. Brock, and J.C. Scott. 1997. Appl Phys Lett 70:2067. 116. Higgins, R.W.T., N.A. Zaidi, and A.P. Monkman. 2001. Adv Funct Mater 11:407. 117. (a) Cao, Y., P. Smith, and A. Heeger. 1993. US-Patent 5,232,631; (b) Geng, Y., Z. Sun, J. Li, X. Jing, X. Wang, and F. Wang. 1999. Polymer 40:5723. 118. Wessling, B. 2000.Handbook of nanostructured materials and nanotechnology, vol 5, ed. H.S. Nalwa, 501–576. New York: Academic Press.
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119. Wessling, B. 2001. Chem Innov 31:34. 120. (a) Becker, H., A. Bu¨sing, A. Falcou, S. Heun, E. Kluge, A. Parham, P. Sto¨el, H. Spreitzer, K. Treicher, and H. Vestweber. 2001. Proc SPIE; (b) Kreuder, W., D. Lupo, J. Salbeck, H. Schenk, and T. Stehlin. European Patent EP 707020. 121. Posdorfer, J., and B. Wessling. Characterization of polyaniline based polymer light-emitting devices during operation by electrochemical impendane spectroscopy. Poster at ICSM 2004. Wollongong, Australia. 122. (a) Meier, M., S. Karg, and W. Riess. 1997. J Appl Phys 82:1961; (b) Kim, S.H., J.W. Jang, K.W. Lee, C.E. Lee, and S.W. Kim. 2003. Solid State Commun 128:143. 123. Esteghamatian, M., and G. Xu. 1995. Synth Met 75:149. 124. Scherbel, J., P.H. Nguyen, G. Paasch, W. Bru¨tting, and M. Schwoerer. 1998. J Appl Phys 83:5045. 125. Paasch, G., and S. Scheinert. 2001. Synth Met 122:144. 126. Posdorfer, J., B. Werner, B. Wessling, S. Heun, and H. Becker. 2003. Proc SPIE San Diego, 188. 127. Becker, H., S. Heun, K. Treacher, A. Bu¨sing, and A. Falcou. 2002. SID Digest Tech Pap 33:780. 128. Boukamp, B.A. 1986. Solid State Ionics 20:31. 129. van Dijken, A., A. Perro, E.A. Meulenkamp, and K. Brunner. 2003. Org Electron 4:131. 130. van Woudenbergh, T., P.W.M. Blom, and J.H. Huiberts. 2003. Appl Phys Lett 82:985. 131. Bru¨tting, W., M. Meier, M. Herold, S. Karg, and M. Schwoerer. 1998. Chem Phys 227:243. 132. Ono, R., M. Kiy, I. Biaggio, and P. Gu¨nter. 2001. Mater Sci Eng B 85:144. 133. Wessling, B. 2001. Chem Innov 31:34. 134. www.ormecon.de 135. Wessling, B. Dispersion–the key tool for understanding and using conductive polymers: Organic nanometals, Lecture at ISCM 2004, Pittsburgh, USA, forthcoming. 136. Wessling, B. 1997.Handbook of organic conductive molecules and polymers, vol. 3, 497–632. New York: Wiley. 137. Wessling, B., D. Srinivasan, G. Rangarajan, T. Mietzner, and W. Lennartz. 2000. Eur Phys J E 2:207. 138. Neural Network Software BrainMaker (California Scientific).
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2 Conducting Polymer Fiber Production and Applications 2.1 2.2
Introduction......................................................................... 2-1 Synthesis of High-Molecular-Weight Polyaniline For Fiber Spinning .............................................................. 2-4 Introduction . Conventional Route for Synthesizing High-Molecular-Weight Polyaniline . Improved Route for Synthesizing High-Molecular-Weight Polyaniline
2.3
Base-Processing Route for the Fabrication of Polyaniline Fibers .............................................................. 2-10 Introduction . Concentrated Emeraldine Base Solutions Containing Secondary Amine Gel Inhibitors . Recent Advances in the Production of Base-Processed Textile Fibers . Asymmetric Polyaniline Hollow Fibers for Membrane Applications
2.4
Acid-Processing Route for the Fabrication of Polyaniline Fibers .............................................................. 2-30 Introduction . Recent Advances in the Production of Acid-Processed Polyaniline Textile Fibers
2.5
Other Methods for Fabricating of Inherently Conducting Polymer Fibers and Fabrics ......................... 2-43 Polyaniline Blend Fibers and Polyaniline
2.6
.
Fabrics Coated with Polypyrrole
Applications of Electrically Conductive Inherently Conducting Polymer Fibers and Fabrics ......................... 2-46 Electrochemical Actuators . Vapor and Humidity Sensors Strain Sensors . Resistive Heaters . Electromagnetic Interference Shielding
2.7
Ian D. Norris and Benjamin R. Mattes
2.1
.
Electrospun Fibers of Inherently Conducting Polymers ... 2-53 Electrospinning Process . Electrospun Polyaniline Fibers . Electrospun MEH–PPV Fibers . Devices Fabricated from Conducting Polymer Electrospun Fibers
2.8
Concluding Remarks......................................................... 2-63
Introduction
The many different types and forms of textile materials that are currently commercially available have one thing in common. The fabrics are passive, and cannot respond or interact by active human control with the environment into which they are placed. However, over the past decade there has been growing 2-1
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interest in the textile industry to integrate electronic circuitry into textiles to modify the functionality of the apparel. In this emerging field of ‘‘smart fabrics and intelligent textiles,’’ a new class of apparel is envisioned, which has active functions embedded in the garment in addition to the traditional properties of clothing. These novel functions or properties are obtained by using special textile fabrics that are interfaced with electronic devices integrated into the garment. Opportunities exist for the use of these smart fabrics and intelligent textiles in the areas of fashion and industrial apparel, residential and commercial interior, military and medical textile markets. For example, applications that have been envisaged for these materials include wearable computing fabric, antistatic garments, garments with electromagnetic shielding capabilities, remote monitoring of a patient’s physiologic status, thermal regulation of sports or military garments, data transfer within clothing for military applications, and sensory fabrics that can dynamically interact with changing environmental conditions and respond to that change in a controllable manner [1–5]. Many of the wires used in the current generation garments with embedded electronic circuitry are cumbersome and awkward. They are simply strapped to the outside of the garments, or carried on the body. This adds bulk and weight to the garment that makes them uncomfortable and impractical for daily use. To enable the integration of electronic circuitry into textiles, a need has arisen for the development of electrically conductive textile fibers and yarns. Electrically conductive textile fabrics have been fabricated from metals such as brass, stainless steel, aluminum, copper, and nickel. These metal fibers have diameters ranging from 1 to 80 mm. Although highly conductive, they are heavier than most textile fibers. One potential solution to overcome these problems has been to coat textile fibers with metals. Currently, synthetic textile fibers that have been coated with silver, copper, and nickel are commercially available. Compared to the homogenous metallic fibers, these metal-clad fibers have a lower overall conductivity of the fiber. However, the most significant disadvantage of these metal-clad fibers is that they cannot be laundered since this coating process reduces the conductivity of the fiber [2]. The metallic-based fibers, either as monofilaments or staple fibers, are often combined with nonconductive fibers to create yarns that can then be stitched, woven, or knitted into electrically conductive fabrics. One of the most cost-effective approaches for producing electrically conductive fibers is the incorporation of conductive fillers, particularly carbon black. However, the incorporation of a sufficient amount of carbon to reach the percolation threshold, sometimes up to 40 wt%, significantly decreases the mechanical properties of the carbon-loaded fibers. Carbon-loaded nylon and polyester fibers available from most commercial fiber producers are generally fabricated using a filled carbon load polymer core or as a sheath of the fiber in order to minimize the decrease in the mechanical properties of the fiber. Although considerably stiffer and more brittle than synthetic textile fibers, graphite fibers can also be incorporated into yarns and fabrics. Composite structures based on graphite fibers, yarns, and fabrics are widely used in the aerospace industry. The academic and commercial industries working in the field of Smart Fabrics and Intelligent Textiles agree that of the few types of conductive fibers that are currently commercially available, carbon-loaded and graphite fibers, metal-clad fibers, and metal fibers, none can fulfill all of the demands envisaged for the applications listed above [1–5]. Textiles based on inherently conducting polymers (ICP) or p-conjugated polymers provide a unique alternative to the range of electrically conductive fibers that are commercially available but there has been very limited commercialization of the electrically conductive ICP fibers and textiles. The major breakthrough in the field of ICP was the discovery in 1977 by Shirakawa and coworkers [6,7] that the partial oxidation of polyacetylene films with gaseous bromine or iodine resulted in a dramatic increase in the electrical conductivity of the polymer with values in the metallic region. Although polyacetylene is the most conductive (up to 105 S=cm) conducting polymer, its poor environmental stability has limited its commercial potential. ICPs have been intensively studied for nearly three decades by both academic and industrial groups around the world both because of their unique physical properties and for their potentially useful application in a variety of exciting technology areas. Initial enthusiasm for these materials in large-scale commercial applications however somewhat waned by the mid-1980s because it was believed that ICPs were largely intractable from a conventional
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polymer-processing viewpoint. The difficulty in processing ICPs is that most cannot be melt-processed, since they decompose at temperatures lower than their melting point. The other difficulty in processing ICPs is that they exhibit poor solubility. Their insolubility arises from the very p-conjugated backbone, which gives rise to their unique electrical properties but which also results in the adoption of a rigid polymer backbone. Fiber spinning can be generally divided into two different processing routes: melt spinning and solution spinning. As the majority of ICPs cannot be melt-processed, the work that has been devoted to the processing of ICP fibers has primarily focused on solution spinning. Effort to fabricate ICP fibers and fabrics using solution-spinning techniques is only now reaching the stage where these materials can be processed using industrially feasible processes. This chapter is primarily devoted to reviewing the advances in the processing of ICP fibers and textiles with a particular emphasis on the processing of polyaniline (PANI) fibers. This chapter also discusses the formation of ICP fibers with submicron dimensions based on an electrostatic-spinning process. This novel spinning process has received increasing attention in the last decade, since the fiber diameters are at least one to three orders of magnitude smaller than fibers made by conventional solution-spinning techniques. Three principal strategies have been developed in order to prepare ICP fibers, regardless of their fiber diameter, which are: 1. Stable homogenous solutions of a highly concentrated (10–30 wt%) ICP polymer is extruded through a spinneret into a coagulation bath that contains a nonsolvent for the polymer, which consequently precipitates the polymer into a fibrous material. This method yields homogenous ICP fibers with good mechanical properties and the highest electrical conductivity. 2. Stable heterogeneous solutions of low concentrations of ICP blended with an insulating polymer and formed into fiber as described above. This method delivers ICP fiber blends that possess the mechanical properties of the insulating polymer, but with electrical conductivity in the semiconducting range. 3. A fiber or textile fabric prepared from insulating fiber materials is coated with an ICP polymer from a dilute solution or through interfacial polymerization techniques to yield a composite material with semiconducting surface conductivity. In a solution-spinning process, the fibers are formed by extruding a concentrated polymer ‘‘dope’’ solution through a spinneret into a coagulant (nonsolvent for the polymer). A spinneret is a die having one or more holes through which the solution is extruded into the coagulation bath. Spinnerets used for industrial fiber spinning typically have between 50 and 200,000 holes with diameters ranging from 10 to 1000 mm. The shape of the spinneret orifice dictates the geometrical shape of the fiber cross section in the solid state. Formation of the fiber occurs rapidly as the dope solution enters the coagulation bath and contacts the nonsolvent, which extracts or withdraws the spinning solvent from the jets of the dope solution exiting from each spinneret hole. This causes the polymer to supersaturate in the dope solution and it then precipitates. As the solvent from the dope solution diffuses from the forming fibers in the coagulation bath, the polymer continues to precipitate and to form a semisolid fiber. While still in the coagulation bath, the fiber achieves sufficient cohesion and strength to remain unbroken upon removal from the bath. The coagulation rate, concentration, and temperature of the spinning solution, and composition, concentration, and temperature of the coagulating liquid are process parameters that affect the final properties of the fiber. The fibers are either continuously removed from the coagulation bath by a take-up or pick-up roll, or godet, followed by a series of further processing operations, including immersion in a series of coagulation baths, washing, wet stretching or orientational drawing, drying, and optionally, hot stretching and annealing to produce selected physical properties for use in textile materials. This further processing of the fiber leads to greater uniformity and microscopic orientation, and hence to better tensile properties, such as high modulus and tensile strength. The processing of soluble ICPs into fibers has largely been limited by the fact that concentrated solutions of these materials are difficult to prepare due to their low solubility and rapid gelation of the concentrated ICP solutions. The dope solution must possess stable rheological properties to produce a fiber with consistent mechanical and, in the case for ICPs, electrical properties. The advances in spinning
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polyaniline fibers have largely been focused on developing a high-molecular-weight polyaniline, methods for preparing concentrated PANI solutions that do not undergo rapid gelation, and the processing conditions, which maximize the fiber’s electrical and mechanical properties. The advances in these three related areas that govern the electrical and mechanical properties of the polyaniline fibers are discussed in this chapter.
2.2 2.2.1
Synthesis of High-Molecular-Weight Polyaniline for Fiber Spinning Introduction
In contrast to most ICPs that cannot be readily processed into fibers due to their intractability, polyaniline can be readily spun into fibers due to its solubility in organic solvents. The history of the polymer now known as polyaniline can be traced to the studies of the dye aniline black. The English dye chemists Green and Woodhead [8,9], working in the British textile industry at the beginning of the last century, proposed an octamer chemical structure for aniline black with five principal oxidation states. Furthermore, these studies found that aniline black was soluble in acetic acid, sulfuric acid, formic acid, and pyridine. The electrical conductivity of polyaniline was not reported until the mid-1980s by MacDiarmid and coworkers [10–13]. It was also shown in these studies that polyaniline actually has three distinct oxidation states and that all other apparent oxidation states are really a mixture of two different oxidation states. Figure 2.1 represents the simplest expression for the polyaniline structure, expressed as a function of its oxidation state consisting of alternating oxidized (ly) and reduced (y) repeat units. When the oxidation state of the polyaniline is fully reduced (y ¼ l), it is called leucoemeraldine; when it is half-oxidized (y ¼ 0.5), it is named emeraldine base; and when it is fully oxidized (y ¼ 0), it is termed pernigraniline. Emeraldine base is a weak basic polymer as there are four nitrogen atoms in the polymer repeat unit, two secondary amines, and two tertiary imines (pKb ¼ 8.6). Polyaniline that is prepared in the emeraldine oxidation state is the composition most often encountered in the literature; and from a technological point of view, it is the most valued since it is environmentally stable and has the most conductive state. Polyaniline has a conduction mechanism that is unique among conducting polymers in that its most highly conducting doped form can be reached by two different processes: protonic acid doping of the emeraldine base (y ¼ 0.5) or oxidative p-type doping of leucoemeraldine (y ¼ 1). Doping of other p-type conducting polymers, e.g., polypyrrole, polyacetylene, and polythiophene results in the formation of a carbonium ion, but the doping of polyaniline results in the formation of a nitrogen salt. This acid–base reaction does not change the oxidation state of the polymer as is the case with redox doping. Upon acid doping, the fully doped form polyaniline is assumed to first form a diamagnetic ‘‘bipolaron’’ form of protonated polyaniline, but magnetic studies of emeraldine salts have shown that they are paramagnetic, which indicates that the bipolarons formed undergo dissociation and protonation to give a delocalized polaron lattice with a polysemiquinone radical cation salt [13]. It can be seen from the alternant resonance forms where the charge and spin are placed on the other nitrogen atoms so that the overall structure is expected to have extensive spin and charge delocalization, which results in a half-filled polaron band (see Figure 2.1). Fully doped emeraldine salt powders have conductivities in the range 1–10 S=cm whereas doped films are more conductive (10–400 S=cm) at room temperature.
NH
NH
N
y Reduced units
FIGURE 2.1
Chemical structure for polyaniline.
Oxidized units
N 1-y x
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In order to generate high-quality fibers possessing good mechanical properties, concentrations of a particular polymer in solution should be in the 10–30 wt% range. Moreover, it is desirable to use the highest molecular weight polymers that will dissolve in solvents in the target concentration range. Tensile strength, modulus, flex life, and impact strength all increase with increasing molecular weight. For dry–wet or wet-fiber-spinning processes that produce high-quality fibers typically from concentrated polyaniline solutions, it is preferable that polyaniline has a weight average molecular weight (Mw) > 120,000 g=mol and number average molecular weight (Mn) > 30,000 g=mol. Before providing an overview on the synthetic routes to obtain high-molecular-weight polyaniline, it is important to note that this procedure somewhat differs from the traditional synthetic routes, which are performed in most laboratories. The chemical oxidation of aniline in acidic aqueous media can be accomplished using a variety of oxidizing agents, which result in the formation of the green, acid-doped, emeraldine salt form of polyaniline as a precipitate. Although the most commonly used oxidizing agent is ammonium persulfate [10–13], other chemical oxidants successfully used for the polymerization of aniline include potassium dichromate [14], iron chloride [15], and hydrogen peroxide [16]. This oxidation reaction is usually carried out in hydrochloric acid or sulfuric acid at a pH between 0 and 2 at a reaction temperature of 08C. For example, one of the most commonly cited methods for producing polyaniline was published by MacDiarmid et al. [12]. In this procedure, the heterogeneous radical chain polymerization of aniline is carried out at 08C in 1 N aqueous hydrochloric acid (HCl) using dropwise addition of an ammonium persulfate solution, and this leads to the formation of HCl-doped emeraldine salt (ES) form of polyaniline. When this polyaniline salt powder is immersed in an excess of a strong aqueous base, it is deprotonated to yield its neutral emeraldine base (EB) form (see Figure 2.2). To determine the reproducibility of this synthesis, the International Union of Pure and Applied Chemistry (IUPAC) selected eight persons from five different countries to carry out polymerizations of aniline following the same preparation protocols [17]. These reactions were carried out at room temperature, and at 08C–28C in 0.2 M (regular acidity) and 1.0 M (high acidity) aqueous HCl solutions. Stoichiometric persulfate oxidant and aniline monomer ratios were adjusted to 1.25 and this resulted in polymer yields between 90% and 100%. It was found that there was excellent reproducibility in ES and EB products generated by the eight individuals performing the reactions. However, it was reported that (a) the reduction in reaction temperature had no marked effect on the ES conductivity and (b) elemental composition (as determined by combustion elemental analysis) of the produced EB polymers at 08C–28C contained 2.3% chlorine via partial benzene-ring substitution
NH
NH
A−
NH
A−
Polyaniline (emeraldine) salt
N
NH
n
−2nH+A− deprotonation
N
NH
NH n
Polyaniline (emeraldine) base
FIGURE 2.2 Change in the polymer backbone upon dedoping the emeraldine salt form of polyaniline to its neutral emeraldine base form.
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with chlorine, especially at the higher HCl acid concentrations. This chlorine ring substitution has consequently prevented its use in some electronic devices. An example of the adverse effect of chlorine ring substitution is in the application of ES thin films as the hole-injecting layer for organic light-emitting diodes (LEDs) [18]. Most polyaniline investigations have employed materials having a weight average molecular weight (Mw) between 50,000 and 100,000 g=mol, and number average molecular weight (Mn) between 10,000 and 30,000 g=mol, which are produced by similar synthetic procedures [19]. The molecular weight distribution of these polyaniline powders does not meet the minimum requirements for producing textile fibers with good mechanical properties. To meet these needs, new synthetic routes have been developed to produce polyaniline, which has a weight average molecular weight (Mw) > 120,000 g=mol and number average molecular weight (Mn) > 30,000 g=mol.
2.2.2
Conventional Route for Synthesizing High-Molecular-Weight Polyaniline
The conventional route for synthesizing high-molecular-weight polyaniline, which is defined as a polymer possessing Mw > 120,000 g=mol and number average molecular weight (Mn) > 30,000 g=mol, is to oxidize aniline in an acidic medium at reaction temperatures less than 208C [20–23]. It is known that addition of certain salts (preferably lithium chloride) to an aqueous solution of aniline hydrochloride allows the reaction mixture to remain liquid at these reaction temperatures whereas the oxidant (preferably ammonium persulfate) is slowly added to the cooled reaction mixture. The reasons why this synthesis produces a high-molecular-weight polymer are straightforward. Aniline polymerizes by a radical cation mechanism and theory shows that this polymerization reaction occurs more favorably in a reaction medium with a high dielectric constant (water ¼ 80, which is high) and at low temperatures. Addition of salts such as lithium chloride (LiCl) further increases the dielectric constant of the reaction mixture and also prevents the reaction mixture from freezing. As the reaction rate decreases, due to lower temperatures, it is thought that the aniline polymerizes preferentially in a head-to-tail manner through the para-position, which is less sterically hindered than the ortho-position. This results in a more linear structure. An early method for preparing high-molecular-weight polyaniline was published by MacDiarmid and coworkers [20,21]. This method involved reducing the standard reaction temperature between 308C and 408C, and by adding between 1 and 6 M of LiCl to the 1.0 M HCl aqueous reaction mixture, thereby producing high-molecular-weight EB. Both increasing the concentration of LiCl in the reaction solution as well as lowering the reaction temperature was shown to increase the weight average molecular weight of the resulting polyaniline, which was found to vary from 250,000 g=mol to greater than 400,000 g=mol by controlling the initial concentration of the reactants. Maintaining the molar ratio of ammonium persulfate to aniline monomer constant while diluting their concentration in the HCl was found to increase the molecular weight of the resulting polymer. However, it was noted that the high-molecular-weight polyanilines produced in accordance with this method exhibited poor solubility and had short gelation times. Acid doping, followed by dedoping with aqueous base, was found to improve solubility in N-methyl-2-pyrrolidinone (NMP). It was speculated that this was likely due to the base catalyzed hydrolysis of the initially long polymer chains into shorter units. Subsequent work by Adams and coworkers [22,23] found that the optimum polymerization temperature for aniline in HCl=LiCl aqueous reaction mixture was approximately 258C, if sufficient ammonium persulfate oxidant is added to polymerize all of the aniline. The resulting weight average molecular weight is about 150,000 g=mol in about 95% yield. If LiCl is added to the oxidant solution as well as to the aniline, the temperature can be reduced to 408C, and only sufficient oxidant may be added to polymerize 40% of the aniline hydrochloride. Moreover, this gives a polymer having a weight average molecular weight of about 250,000 g=mol. If additional oxidant is added, the oxidant reacts with the polyaniline as well as the monomer, giving lower molecular weight material. It was furthermore noted that the 13C NMR spectra of
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the resulting higher molecular weight polyaniline contained fewer defect sites than the material synthesized at higher temperatures using ammonium persulfate as the oxidant. Defects in the polyaniline backbone are defined as any structural deformation of the polyaniline linear chain that disrupts the conjugation of alternating single and double bonds, e.g., chain branching and cross-linking.
2.2.3 Improved Route for Synthesizing High-Molecular-Weight Polyaniline As noted earlier, if aniline is polymerized in an acid reaction mixture with large amounts of LiCl present, especially if the acid is HCl, significant ring chlorination occurs (typically 1% by weight of the base polymer is ring-bound chlorine through covalent bond formation) [17]. Moreover, the addition of large amounts of LiCl to the reaction mixture greatly increases the final costs of the polymer because (a) LiCl is an expensive additive and (b) it is difficult to separate from the remaining aqueous HCl reaction mixture, thereby increasing the costs associated with hazardous waste removal. Therefore from a production standpoint, it is advantageous to eliminate the use of LiCl altogether. Therefore, the research that was performed at Santa Fe Science and Technology (SFST) was to look at alternatives to using a HCl=LiCl reaction mixture for synthesizing high-molecular-weight polyaniline. The main reasons for doing this research are: 1. To eliminate ring-substituted chlorine, found after polymerization in HCl 2. To find a cheaper way to polymerize aniline at low temperatures 3. To synthesize defect-free polyaniline with Mw > 120,000 g=mol We have shown in our laboratory that the HCl=LiCl reaction media could be replaced with other commercially viable inorganic acids, such as highly concentrated solutions of sulfuric acid and phosphoric acid, and organic acids, such as highly concentrated solutions of formic acid, acetic acid, and difluoroacetic acid [24]. By carefully controlling the water content of these acids, it is possible to keep the reaction mixture fluid and thus polymerize aniline at temperatures as low as 508C. To obtain highmolecular-weight polyaniline suitable for fiber spinning, batch reactions were performed between 08C and 508C by adding an ammonium persulfate solution at a chosen rate to a cold mixture of aniline and a suitable acid. Acid concentrations and types were chosen such that the reaction mixture remained fluid at the low reaction temperatures whereas the resulting polymer was not degraded by the presence of the high acid concentration. Important processing variables, which we have found to affect the molecular weight of polyaniline include: (a) reaction temperature, (b) total reaction time, (c) aniline monomer concentration, (d) choice of acids, (e) acid concentration, (f) conversion of aniline monomer, (g) amount of oxidant added, and (h) oxidant addition rate. Later on in this section, we will primarily focus on the effect of reaction temperature on the molecular weight of polyaniline synthesized in highly concentrated solutions of sulfuric acid or phosphoric acid. The preponderance of patent or scientific literature regarding polyaniline synthesis in aqueous media reports synthetic conditions whereby the concentration of the acid is measurable on the pH scale and the acid most frequently reported is HCl. The activity and concentration of the hydronium ion are obtained by measurements of pH by ion-selective electrodes or pH paper–containing indicators. However, such pH measurements are valid only in single-solvent systems, typically water, for very dilute concentrations of an acid. However for very concentrated acid solutions, such as those used to polymerize aniline at subzero temperatures, the preferred measure of the ability of an acid to dissociate a proton from an indicator, according to HBþ $ HþþB, is the Hammett acidity function, H0 given by H0 ¼ pKHBþ log
CHBþ CB
(2:1)
where CHBþ and CB are the concentrations of the two forms of a protonated and nonprotonated indicator, respectively, in an equilibrium mixture [25]. Indicator compounds used to determine the Hammett acidity function (H0) include aniline, or more commonly substituted anilines such as
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p-nitroaniline. Once H0 is determined, Equation 2.1 can be used directly, in a similar manner to pH, to obtain unknown acidity constants from ionization ratio measurements, that is
280 270 Freezing point (K)
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260
pKHBþ ¼ Ho þ log [CHBþ =CB ]
(2:2)
250
Concentrations of HBþ and B are measurable by spectroscopy, and pKa values of the acids HBþ are 240 well known. Hammett acidity function, H0, scales are useful for comparing different acid media for 230 acid strength. Typically, the Hammett acidity 220 functions for the reaction mixtures to synthesize 0 5 10 15 20 25 30 35 40 high-molecular-weight polyaniline are in the Sulfuric acid concentration (wt%) range: 2 H0 0.5. Figure 2.3 shows the freezing point for concenFIGURE 2.3 Freezing point depression of water as a trated aqueous sulfuric acid solutions, from which function of increasing sulfuric acid concentration. it can be seen that using sulfuric acid solutions the reaction medium allows temperatures as low as 233 K (408C) to be attained without the reaction mixture freezing, and without the use of freezing point lowering salts such as LiCl. Using the data in Figure 2.3, different syntheses of polyaniline reactions were performed at different temperatures in a range from 08C to 458C, using the amount of sulfuric acid necessary to just prevent the mixture freezing at that temperature. Similar to what has been previously been published, the stoichiometric persulfate oxidant=aniline monomer ratio was 1.25:1 and this enabled polymer yields greater than 90% to be successfully obtained for all syntheses. The amount of aniline monomer added to the polymerization reaction mixture was chosen to be between 0.3 and 2.0 M. After allowing the reaction to proceed for 24 h, the contents of the reaction vessel were filtered, and washed with water until a colorless filtrate was obtained. The filter cake was subsequently deprotonated using a 2 wt% ammonium hydroxide solution, the suspension was refiltered, rewashed (with a final wash of 2-propanol), and then dried under vacuum. The molecular weights of these polyaniline (emeraldine base) powders were determined by gel permeation chromatography (GPC) using polystyrene standards of different molecular weights. The polyaniline is dissolved in the polar aprotic solvent N-methyl-2-Pyrrolidinone (NMP). An ionic salt is added to prevent aggregation of the polyaniline chains; otherwise, a non-Gaussian molecular weight distribution is observed [26]. Typically, ionic salts used to deaggregate polyaniline include lithium and ammonium salts, such as lithium chloride, lithium bromide, lithium tetrafluoroborate, and lithium formate. With the ionic salt lithium chloride dissolved in the eluent, the polyanilines exhibited single-peak gel-permeating chromatograms. Table 2.1 shows the molecular weight distribution for the 10 different syntheses of TABLE 2.1 Effect of Reaction Temperature on the Molecular Weight Distribution of Polyaniline Synthesized in Concentrated Sulfuric Acid Aqueous Solutions Reaction Temperature (8C) 0.0 10.0 15.0 25.0 30.0 35.0 37.5 40.0 42.5 45.0
H2SO4 (wt%)
H0 at the Reaction Temperature
Mp (g=mol)
Mw (g=mol)
Mn (g=mol)
4.5 16.9 21.0 26.2 28.2 29.8 30.5 31.1 31.5 32.0
0.13 0.80 1.08 1.44 1.55 1.72 1.76 1.81 1.84 1.87
67,800 91,200 107,000 155,000 200,000 229,000 207,000 227,000 219,000 233,000
77,500 111,000 148,000 178,000 235,000 278,000 246,000 366,000 269,000 307,000
25,700 30,500 38,900 56,400 64,000 65,500 56,400 65,800 69,000 67,000
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polyaniline. As expected, the results in Table 2.1 show a gradual increase in molecular weight (both Mw and Mn) with decreasing reaction 270 temperature and Hammett acidity function. Similarly, Figure 2.4 shows the freezing 260 point for concentrated aqueous phosphoric acid solutions and illustrates that using con250 centrated phosphoric acid solutions as the reaction medium allows for the polymeriza240 tion of aniline at temperatures as low as 223 K (558C) without the use of ionic salt addi230 tives. A series of reactions similar to those in sulfuric acid at different temperatures between 220 108C and 558C were carried out in 60 wt% 0 10 20 30 40 50 60 phosphoric acid reaction mixture. The stoiPhosphoric acid concentration (wt%) chiometric persulfate oxidant=aniline monoFIGURE 2.4 Freezing point depression of water as a mer ratio was again 1.25:1 and this enabled function of increasing phosphoric acid concentration. polymer yields greater than 90% to be successfully obtained for all syntheses. The purpose was to show that a ‘‘standard’’ reaction mixture and procedure could be utilized, where the only variables are the temperature and time of reaction. As noted earlier, different temperatures result in different molecular weight polyaniline. The total reaction time may be varied to suit the rate of reaction at a particular temperature. The Hammett acidity of 60% phosphoric acid is about 1.6, at 208C, and increases with decreasing temperature. The total reaction time was between 43 and 46 h, with the exception of the polyaniline powder synthesized at 558C, which was 90 h due to the slower reaction kinetics. The molecular weight values were measured by GPC using the same procedure as described for the polyaniline powders listed in Table 2.1. The molecular weights, especially the weight-average molecular weights, were higher than those obtained in sulfuric acid at the same reaction temperature. This could be attributed to the improved stirring that resulted from performing the reactions in concentrated phosphoric acid solutions (Table 2.2). Alternately, since sulfuric acid is a stronger acid than phosphoric acid, it could possibly cause excess hydrolysis and degradation of polymer backbone. It should be noted that the reactions below 458C resulted in lower molecular weights. For producing high-molecularweight polyaniline for fiber spinning, it was found that the optimal synthetic conditions were a reaction temperature of 358C in a 60 wt% phosphoric acid, since polyaniline synthesized at lower temperatures does not increase either the weight or number average molecular weights. Finally, elemental analysis of EB powders synthesized in phosphoric acid showed no evidence for the presence of chlorine, which Freezing point (K)
280
TABLE 2.2 Effect of Reaction Temperature on the Molecular Weight Distribution of Polyaniline Synthesized in 60 wt% Phosphoric Acid Aqueous Solution Reaction Temperature (8C) 10 20 25 30 35 40 45 50 55
Mp (g=mol)
Mw (g=mol)
Mn (g=mol)
159,000 157,000 206,000 203,000 206,000 130,000 136,000 83,600 87,500
302,000 296,000 403,000 475,000 426,000 273,000 267,000 145,000 158,000
29,000 29,300 50,000 52,400 75,500 33,700 34,600 17,400 21,500
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indicates that ring chlorination was not present in the polyaniline powders. This is in contrast to the synthesis of polyaniline at 258C in our laboratory using a 1.0 M HCl þ 6.0 M LiCl reaction mixture that contained 2%–3% chlorine due to partial benzene ring substitution with chlorine.
2.3 2.3.1
Base-Processing Route for the Fabrication of Polyaniline Fibers Introduction
Processing of polyaniline into fibers from its neutral, emeraldine base form is highly attractive since the polymer is not a rigid rod, and therefore fiber spinning should be possible once an appropriate solvent for the polymer is found. Angelopoulos et al. [27] first demonstrated that EB is soluble in a range of organic solvents including NMP, acetic acid, formic acid, dimethyl sulfoxide, and dimethyl formamide. The dissolution of emeraldine base in NMP has been the most widely practiced route for processing polyaniline into both films and fibers. Subsequent work demonstrated that EB solutions in NMP at polymer concentrations in excess of 5–6 wt% formed gels after a short period of time [19,28,29]. The spinning of emeraldine base fibers from NMP solutions was first reported by MacDiarmid and coworkers [30,31]. These fibers were made from concentrated (20 wt%) EB (synthesized at 08C) solutions dissolved in NMP. The tensile strength of the fibers after stretching 3–4 times its original length was 366 MPa. Although these fibers exhibit only modest mechanical properties in comparison with traditional high-strength textile fibers, this preliminary study demonstrated that fiber spinning from concentrated EB solutions is possible. Not surprisingly, rapid gelation of these concentrated polyaniline EB solutions was observed. The formation of these polyaniline gels at concentrations lower than that required for fiber spinning necessitates the use of a binary solvent system for concentrated EB=NMP solutions. Cohen and coworkers [32,33] demonstrated that addition of certain amines as cosolvents, such as pyrrolidine to concentrated EB=NMP, or dissolving EB in certain amine solvents, such as 1,4-diaminocyclohexane, delayed the onset gelation of these solutions thus allowing for fiber spinning on a meaningful scale. It was postulated that the gelation delay of the EB solution was related to the disruption of microcrystalline regions in the solution by the amine cosolvent. These studies revealed that continuous dry-jet wet spinning could be carried out from concentrated EB=1,4-diaminocyclohexane solutions (10–20 wt% EB) when using 1,4-diaminocyclohexane, with the EB fiber possessing an initial tensile strength of 80 MPa or 0.8 g per denier (gpd).* It was found that the as-spun EB fiber was crystalline and could be oriented at high temperatures, e.g., 2158C. The drawn fibers have an average tensile strength of 400 MPa (3.9 gpd) with a modulus of 8.5 GPa (83 gpd). The highest electrical conductivities obtained with drawn fibers doped with aqueous sulfuric acid and hydrochloric acid were 321 and 158 S=cm, respectively. No cause was given for the difference in the conductivity of these fibers. However, doping the fiber with HCl caused the tensile strength of the doped fiber to be reduced by 64% to 1.4 gpd. Gregory and coworkers [29,34] studied the rheological properties of concentrated polyaniline solutions (8–10 wt% EB) dissolved in NMP, 0.5 wt% LiCl dissolved in NMP and dimethylpropylene urea (DMPU). The addition of salts such as LiCl to fiber-spinning solutions is a commonly used technique for reducing the aggregation of polymers in solution [35]. It was found that the addition of LiCl to the NMP solvent delayed the onset of gelation relative to the NMP alone, but did not provide a sufficiently stable window to enable fiber production. It was found, however, that the gelation time and solution stability of polyaniline in DMPU is greatly increased. For example, a 10 wt% EB solution in NMP gelled within 60 min, whereas a 17.5 wt% EB solution in DMPU did not show any significant change in its *Tensile strengths (and modulus) are usually reported in the textile industry in the units of grams per denier (gpd—breaking force in grams per 9000 m of filament) whereas the general scientific community uses SI units. Where possible, both sets of units are listed. To convert from gpd to SI units the density of the fiber required, which is often not reported in these studies.
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viscosity for as long as 400 min. This indicates that DMPU is a more suitable solvent than NMP for wetspinning polyaniline fibers in their EB form. Jain and Gregory [34] later showed that polyaniline fiber could be obtained from a 15 wt% low-molecular-weight EB=DMPU solution into a NMP=water coagulation bath. The tensile strength, elongation, and modulus of the as-spun polyaniline fiber processed from DMPU were found to be 1.5–3.0 gpd, 12%–18%, and 8–12 gpd, respectively. Upon doping the as-spun fiber that had been stretched four times its original length with methanesulfonic acid dissolved in acetic acid, the conductivity of the fiber was found to be 350 S=cm. Gregory and Eaiprasersaks [36] later showed that it was possible to spin fiber from a highly concentrated leucoemeraldine base solution dissolved in DMPU. Since polyaniline is fully reduced in its leucoemeraldine base form, it only contains amine nitrogen groups along the polymer chain. The elimination of imine nitrogen groups from the polymer chain reduces the polyaniline’s ability to form secondary hydrogen bonds with neighboring chains. The formation of hydrogen bonds between the amine and imine nitrogen groups leads to the gelation of concentrated EB solutions. The major disadvantage of this approach is that the as-spun leucoemeraldine fiber must be reoxidized to its emeraldine oxidation state in order for the fiber to become electrically conductive. The as-spun fiber could be obtained from a 15 wt% leucoemeraldine=DMPU solution into an NMP=water coagulation bath having a tensile strength, elongation, and modulus of 3.6 gpd, 15%, and 89 gpd, respectively. The electrical conductivity of the fiber after oxidation and doping yield a conductivity of 150 S=cm for the doped fiber. After doping the fiber tensile strength, elongation, and modulus of 1.9 gpd, 23%, and 41 gpd, respectively. The mechanical properties were superior to the doped fiber that was processed from the concentrated EB=DMPU solution, and are strong enough to be processed using conventional textileprocessing equipment. It was observed by Hsu and Epstein [37] that solutions of the EB form of poly(o-toluidine) dissolved in NMP did not show any signs of gelation at the same polymer concentrations at which the parent polyaniline undergoes rapid gelation. The authors speculated that this difference in solution behavior could not be simply attributed to the lower molecular weight of poly(o-toluidine), but rather the methyl groups on the aromatic rings prevented poly(o-toluidine) from crystallizing. It was also shown that poly(o-toluidine) fibers could be spun from a 33 wt% solution of the EB form of poly(o-toluidine) dissolved in NMP by extruding the solution into a water coagulation bath. The tensile strength, elongation, and modulus of the as-spun poly(o-toluidine) fibers were found to be 0.8 gpd, 48.0%, and 18 gpd, respectively. The tensile strength of the as-spun fiber increased to 2.4 gpd with a concurrent decrease of elongation from 48% to 17.0% and an increase of modulus from 18 to 40 gpd upon stretching the fiber three times over a hot pin at 2208C. This process also resulted in the alignment of the polymer chains. The alignment enhanced electrical conductivity of the HCl-doped fiber from 0.02 to 0.1 S=cm. Mattes and coworkers [38,39] reported on the fabrication of EB fibers from concentrated (20 wt%) high-molecular-weight (Mw ¼ 135,000 g=mol) EB solutions, which were stabilized by small amounts of the secondary amine, 2-methylaziridine, into the EB=NMP solution. These solutions had gelation times that varied from a few hours to several days depending on the molar ratio of 2-methylaziridine to EB tetramer-repeating unit in the NMP solvent. The as-spun fibers were amenable to thermal stretching with maximum draw ratios of 4:1. The four-times-stretched fiber exhibits a higher Young’s modulus and yield strength (1.85 GPa, 50 MPa) than that of the as-spun fiber (540 MPa, 12 MPa). In contrast the four-times-stretched fiber breaks at a strain of 6%, which is about 30% smaller than the failure strain of as-spun fiber (9%). This difference in mechanical properties is in all likelihood due to the increased density and chain alignment in the four-times-stretched fiber. The measured densities of the as-spun and four-times-stretched fiber are 0.52 and 0.92 g=cm3, respectively. This suggests that thermal stretching increases the Young’s modulus, yielding strength, and failure strength of the fiber. It was found that doping the fiber with an inorganic acid (HCl) embrittled the fiber whereas doping with organic acids (acetic acid, benzenephosphinic acid) did not appreciably affect the mechanical strength. For fibers with a four times draw ratio, the benzenephosphinic acid-doped fiber exhibited the highest conductivity with a value of 10.3 S=cm. The fiber conductivity when doped with the acetic acid and HCl was (0.7 and 3.4 S=cm, respectively). Despite giving good mechanical properties, the secondary
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amine used to process these fibers, 2-methylaziridine, is highly toxic, which inhibits its use for commercial scale fiber spinning.
2.3.2
Concentrated Emeraldine Base Solutions Containing Secondary Amine Gel Inhibitors
As noted earlier, although NMP is commonly used to process polyaniline in its EB form, rapid gelation occurs during short periods of time when the EB concentration exceeds 5 wt%. Below 5 wt%, interchain H-bonding is weak and solutions remain stable for several hours without gelling. However, if the EB concentration is increased beyond 5 wt%, irreversible aggregation and the development of a strong, physically cross-linked three-dimensional gel network occurs in a short period of time because the polymer amine nitrogens, which are not associated with solvent molecules reform their interchain H-bonds with the nearest neighboring imine nitrogens. It is important to note that the imine nitrogens (tertiary amines) cannot participate in any hydrogen-bonding interaction with amide solvent molecules such as NMP, since the latter has no protons to donate to a hydrogen bond with the polymer. Mattes and coworkers [38–44] showed that concentrated emeraldine base solutions could be stabilized for extended periods of time through the addition of trace amounts of secondary amines into the EB=NMP solution. They coined the term ‘‘gel inhibitors’’ to describe the function of these secondary amines additives since they (a) assist in achieving complete dissolution of the polymer, (b) decrease overall solution viscosity compared with initial solution viscosity, and (c) greatly prolong the shelf life of the polymer solutions by increasing time to gelation. The mechanism postulated for this phenomenon is related to the formation of a hydrogen bond adduct between the proton of the secondary amine additive and the lone pair of electrons of the imine nitrogen repeat unit. Therefore, the secondary amine additive and NMP impede the EB nitrogen from inter- and intrachain hydrogen bonding. For example, hydrogen-bonding formation among NMP, 2-methylaziridine (2MA), and EB molecules are schematically illustrated in Figure 2.5. Secondary amines other than 2-methlaziridine form similar hydrogenbond networks with the polyaniline chain. It was shown that concentrated solutions of EB will dissolve in polar aprotic solvents such as NMP when near-stoichiometric amounts of 2MA are mixed with the solvent, i.e., when the molar ratio of 2MA=EB tetramer repeat unit (Mw ¼ 362 g=mol) is near 2 [38–42]. Formation of these solutions requires heat or high-shear mixing. Figure 2.6 illustrates the three distinct rheological regions observed
CH3 N CH3
N
O
H
H
N
N N
N
H
H
N
O
CH3 N CH3
FIGURE 2.5 Schematic diagram suggesting the hydrogen-bond formations between EB, NMP, and 2MA molecules. (Reprinted from Yang, D. and Mattes, B.R., Synth. Met., 101, 746, 1999. With permission from Elsevier.)
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Viscosity
for these solutions. The decrease in initial solution viscosity in Region I is attributed to the energy required for breaking the intermolecular hydrogen bonds in the EB powder and the concurrent formation of the NMP–EB–2MA hydrogen-bond I II III complex. Region II represents the period of time during which there is no change in the viscosity of the system. In Region II, the solution has reached a stable equilibrium state in which approximately two secondary amine additive molecules have complexed with the two imine nitrogens in the Time polymer repeat unit to form a stable hydrogen- FIGURE 2.6 Representative viscosity vs. time curve bond interaction complex. After some time, the showing three distinct regions: (I) 2MA–EB hydrogensolution phase separates due to the prolonged bond complex formation region, (II) stable equilibrium effects of continuous shearing. Region III shows viscosity region, and (III) gelation region. (Reprinted the gelation behavior observed for the reactive- from Yang, D. and Mattes, B.R., Synth. Met., 101, 746, polymer mass. During this period, the intermo- 1999. With permission from Elsevier.) lecular polymer hydrogen bonds are reestablished, which leads to solidification as three-dimensional physically cross-linked network forms. The major drawback of using secondary amines is that some, such as pyrrolidine, can react with the EB chains, especially when the molar ratio of gel inhibitor to EB tetramer repeat unit is greater 4:1, leading to chemical reduction and fragmentation products. For example, Han and Jeng [45] showed that in EB solutions containing pyrrolidine, the polyaniline underwent chemical reaction due to concurrent reduction and nucleophilic substitution of the H-atom on the semiquinone ring with pyrrolidine molecule. In other words, this reaction reduces the emeraldine base into a leucoemeraldine-like derivative through a ring substitution. Using NMR, this aging process could be detected as early as 10 min after the EB=pyrrolidine solutions were prepared. The chemically degraded PANI possesses lower conductivity and poor mechanical properties. Yang and coworkers [43,44] modeled the conformation of a tetramer of aniline (with the terminal amine removed) in the EB form (Figure 2.7) to study the possible interactions between the polymer and the secondary amine. The resulting minimized structure shows the expected twist of the aromatic rings along the chain. The distance between the two ortho-hydrogens, next to the imine nitrogen of
4.53 Å Hydrogen
Carbon
Imine nitrogen
Amine nitrogen
Lone pair electrons N
N H
N
FIGURE 2.7 Computer model of the EB tetramer showing the two ortho-hydrogens around the imine (above), and the structure of the aniline EB (with the terminal amines removed) (below). (Reprinted from Yang, D., Zuccarello, G., and Mattes, B.R., Macromolecules, 35, 5304, 2002. With permission from the American Chemical Society.)
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the EB, was calculated to be 4.53 A˚ (see Figure 2.7). An amine that approaches the imine lone pair electrons in order to form a hydrogen bond has to be small enough to avoid interacting with these two flanking hydrogens. To further elucidate how these secondary amines interact with EB molecules, Yang and coworkers [43,44] investigated several different cyclic secondary amines with different basicity and geometric factors (see Table 2.3). All of these secondary amines behave as gel inhibitors. A 20% solution of EB (Mw of 90,000 g=mol) in NMP=secondary amine with a secondary amine:EB tetramer repeat unit molar ratio of TABLE 2.3 Summary of the Width, Depth, and Molecular Size and pKa of Cyclic Secondary Amines and the Gelation Time Obtained for a 20% Solution of EB (Mw of 90,000 g=mol) in NMP Containing the Secondary Amine Name and Structure 2-Methylaziridine (2MA)
Width (A˚), Depth (A˚), Molecular size (A˚)
pKa
Gelation Time (h)
2.62 2.30 4.37
8.5
>300
2.77 1.87 3.80
11.3
>300
4.08 1.82 4.22
11.5
>300
4.19 3.10 4.90
11.3
>300
4.14 3.09 5.21
11.4
>300
4.05 3.11 5.89
11.2
>300
N H Azetidine (AZ)
NH Pyrrolidine (PY)
N H Piperidine (PP)
N H Hexamethyleneimine (HXMI)
N H Heptamethyleneimine (HPMI)
NH Source: Reprinted from Yang, D., Zuccarello, G., and Mattes, B.R., Macromolecules, 35, 5304, 2002. With permission from the American Chemical Society.
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2.0 did not show any indication of gelation after 300 h. The degradation of the polyaniline chains was monitored using UV–vis spectroscopy to detect changes in the oxidation state of the polyaniline chains and using GPC to detect reduced molecular weights of the polyaniline chains if the aged EB solutions are used. In these studies, the secondary amine to EB tetramer repeat unit molar ratio was 2:1 for all samples. Based on the experimental observations, it was hypothesized that besides serving as proton donors to form hydrogen bonds, the amines can also serve as nucleophilic agents to attack the quinoid ring at different sites, resulting in different products [45]. Figure 2.8 illustrates two possible reactions between an amine and EB. In the first reaction, a secondary amine with a large degree of steric hindrance such as piperidine (PP) attacks the quinoid ring and leads to a ring substitution. This was also proposed by Han and Jeng [45] who hypothesized that protonation of the imine nitrogen could possibly promote nucleophilic attack at the meta-position of the protonated quinoid ring. Therefore, the less stable quinoid ring would be reduced to the more stable benzenoid by ring substitution. This structural change decreases the conductivity of the HCl-doped film but essentially leaves its mechanical properties intact [44]. The second reaction, showing a secondary amine, such as azetidine (AZ), with a small degree of steric hindrance involves direct attack at the }N¼C< site. When moisture or another impurity is present, chain scission together with polymer reduction may take place to produce a substantially lower molecular weight polymer. Matveeva et al. [46] proposed a similar reaction between water and EB molecules, leading to chain scission. Chain scission helps explain the poor mechanical properties observed for the films made from the EB=NMP=AZ and EB=NMP=PY solutions and it also correlates with the 15-fold reductions in the peak molecular weight values of the polyaniline chains. Furthermore, it was observed that the UV–vis spectra more closely resembled that of partially reduced aniline oligomers than that of polyaniline in its EB oxidation state. Therefore, it appears likely that both chain scission and ring substitution occurs in aged EB=NMP=AZ solutions and that this reaction is largely responsible for the change in the UV–vis spectra. It appears that piperidine mainly reduces the EB structure by ring substitution. Regarding other secondary amines, they may interact with EB through either of the two reactions. Based on the mechanical properties, it was deduced that pyrrolidine, due to its smaller degree of steric hindrance,
Reaction 1: Leading to ring substitution and reduction H N N H N
N H
H
N H
N
N
—Q substitution, thus chain scission and reduction Reaction 2: Leading to −N— N
C
H N
N H
+ N
− NH +
C N
H2O NH2 Rearrange
N HO NH or H2O
HO
or N
N OH
FIGURE 2.8 Schematic diagram of sterically different secondary amines attacking an imine center in an EB tetramer repeat unit through different reactions mechanisms. (Reprinted from Yang, D., Zuccarello, G., and Mattes, B.R., Macromolecules, 35, 5304, 2003. With permission from the American Chemical Society.)
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tended to react with EB through the second reaction pathway whereas heptamethyleneimine (HPMI) tended to preferentially react with EB through the first reaction pathway. It was concluded that in the case of secondary amines possessing similar steric hindrance values, e.g., AZ vs. 2MA, it is the basicity of the additive that governs the strength and nature of the chemical interaction or reaction with the PANI. This conclusion is supported by data acquired from both the UV–vis spectral characterization of the elution peaks and the GPC analysis (DMp based on retention times). The amine having weaker basicity has a reduced tendency to react with the polymer and forms more thermodynamically stable solutions at high concentration. Conversely, when the secondary amines possess comparable pKa values, it is the steric hindrance factor which determines the extent and nature of the interaction with the EB imine nitrogen. The amine with the smallest steric hindrance has the greatest tendency to chemically degrade the polymer structure. In summary, the reactivity strength among the cyclic amines toward EB decreases in the following order: AZ > PY > HPMI > HXMI > PP > 2MA. In the search for secondary amines that inhibit gelation of concentrated EB=NMP solutions but do not degrade the EB structure, Yang et al. [43] investigated the gelation inhibition and degradation characteristics of 39 secondary amines. It was found that when both the width and depth of a secondary amine is <4.53 A˚ and its pKa is >7.7, then this amine could extend the gelation times of 20 wt% EB=NMP solutions for more than 12 h (see Figure 2.9). However, when both the width and depth of these amines are >4.53 A˚, these amines neither prolong gelation time nor appreciably degrade EB. It should be emphasized that amines with small width and depth and strong basicity, such as azetidine and pyrrolidine, have a tendency to significantly degrade the EB structures. Based on these studies, secondary amines can therefore be divided into three categories: 1. Nondegrading GIs able to prolong gel time for at least 12 h, with no appreciable degradation of EB structure 2. Degrading GIs able to prolong gel time for at least 12 h, but with appreciable degradation of EB structure 3. Non-GIs unable to prolong gel time for longer than 1 h
10 9
non-GIs GIs
˚) Max {width, depth}(A
8 7 6 5 4 3 2 4
5
6
7
8
9
10
11
12
pKa
FIGURE 2.9 A plot of GI efficiency based on the maximum of width and depth of amine vs. pKa (GI: solution gelation time >12 h; non-GIs: solution gelation time <1 h). (Reprinted from Yang, D., Zuccarello, G., and Mattes, B.R., Macromolecules, 35, 5304, 2004. With permission from the American Chemical Society.)
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2.3.3 Recent Advances in the Production of Base-Processed Textile Fibers Some of the limitations of the previous studies include limits in the length of the EB fiber obtained due to the dope solution short lifetime and nature of the experiment, decreased molecular weight polyaniline (Mw < 100,000 g=mol), and the destructive nature and toxicity of the amine gel inhibitors. For these reasons, investigations have been performed that focused on spinning EB fibers on a pilot plant scale using higher molecular weight polyaniline than has been traditionally used (Mw > 100,000 g=mol). In conjunction, secondary amine gel inhibitors with lower toxicity and reduced propensity to attack the polymer backbone have been investigated. 2.3.3.1 Spinning of Emeraldine Base Fibers Polyaniline solutions for fiber spinning were prepared using the secondary amines HPMI and 4-methylpiperidine (4-MP) as stabilizers for the concentrated EB=NMP solutions. These reduced the EB more slowly than secondary amines such as 2-methylaziridine [43]. Polyaniline fibers were processed from dope solutions in which the concentration of EB base powder (Mw 150,000 g=mol) varied from 17.5 to 25 wt% polyaniline in the case of the HPMI=NMP or 20 wt% of EB base powder (Mw 200,000 g=mol) for the 4-MP. The stoichiometric ratio between the HPMI molecules and the EB tetramer repeat unit required to stabilize the dope solution was 1.1:1 whereas the 4-MP:EB tetramer ratio was 1.2:1. The EB was mechanically stirred into the secondary amine=NMP solution until a smooth, lump-free solution was obtained. Understanding the rheology of the fiber-spinning dope solutions is important to enable the fabrication of the polyaniline fibers with reproducible electrical and mechanical properties. In our studies, a Brookfield RVDV-III cone and plate rheometer was used to measure the viscosity of the concentrated EB solutions. The gelation studies of these solutions were conducted at 258C and a constant shear rate of 0.8 s1 using cone spindle (CP-52) with a semivertical cone angle of 878. Figure 2.10 illustrates the viscosity behavior of the fiber-spinning dope solutions (EB=NMP=HPMI) with different EB concentrations as a
3.0 ⫻ 105 17.5 wt% 20.0 wt% 22.5 wt% 25.0 wt%
Viscosity (cP)
2.5 ⫻ 105
2.0 ⫻ 105
1.5 ⫻ 105
1.0 ⫻ 105
0.5 ⫻ 104
Gelation time
0 0
10
20
30
40
50
60
70
Time (h)
FIGURE 2.10 Rheological behavior of concentrated EB=NMP=HPMI solutions for different EB concentrations. (Reprinted from Yang, D., Fadeev, A.G., Adams, P.N., and Mattes, B.R., Proc. SPIE, 4329, 59, 2001. With permission from the International Society of Optical Engineering.)
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Conjugated Polymers: Processing and Applications TABLE 2.4 Summary of Solution Viscosity in Region II, Gelation Time at 258C for the EB=NMP=HPMI Solutions Made from Different EB Concentrations EB Concentration (wt%) 17.5 20.0 22.5 25.0
Viscosity in Region II (cP)
Gelation Time (h)
4,200 18,000 74,000 135,000
32 20 8.5 6.5
Source: Reprinted from Yang, D., Fadeev, A.G., Adams, P.N., and Mattes, B.R., Proc. SPIE, 4329, 59, 2001. With permission from the International Society of Optical Engineering.
function of time. Three distinct rheological regions were observed for these solutions regardless of the EB concentration and the molecular weight. This observation is similar to what was reported by Yang et al. [42] for concentrated EB=NMP solutions that were stabilized using the secondary amine 2-methylaziridine. Hydrogen-bond formation between HPMI and EB is believed to be responsible for the viscosity decreasing in Region I. In Region II, hydrogen-bond formation between HPMI and EB reaches a metastable equilibrium and the viscosity remains constant. Eventually, inter- and intrapolymer chains hydrogen bonding begins to dominate and the viscosity increases as the solution begins to gel, Region III. Table 2.4 summarizes the viscosity effects of the EB concentration in Region II as well as the gelation times of these concentrated EB solutions. As the EB concentration in the dope solution increases from 17.5 to 25 wt%, the gelation time decreases (from 32 to 8 h) whereas the viscosity in Region II increases (4,200 to 135,000 cP). The 20 wt% EB=NMP=4-NMP dope solution possessed a stable Region II viscosity of 120,000 cP for at least 6 h. Since it is difficult to process gelled material (Region III), the optimal processing zone for these solutions is Region II. The length of time in Region II in which process the polyaniline fiber before the onset of gelation can vary from a few hours to several days, depending on the EB concentration and molecular weight as well as the stoichiometric ratio between gel inhibitor molecules and the EB tetramer repeat unit. From the rheological study of the EB dope solutions, we can predict its processing zone and estimate the force required by the gear pump during spinning. The concentrated EB=NMP=gel inhibitor solution is loaded into the cartridge and pressurized with nitrogen gas at 60 psi to deliver a steady flow rate of the dope solution to the gear pump. The gear pump pushes the dope solution through three stages of filters (230, 140, and 90 mm). The solution is filtered to remove any undissolved or gel particles that could become dislodged in the spinneret orifice. The filtered dope solution is then extruded through a 100 mm diameter single-hole spinneret (l=d of 2) into a 1 m long, temperature-controlled water coagulation bath. In the coagulation bath, the dope solution solidifies and the nascent fiber is continuously taken up on a 16.5 cm diameter stainless steel godet drum equipped with a digital controller for precise take-up speed control. In these studies, the take-up speed was varied between 10 and 20 m=min. The nascent fiber is then passed through two washing godet baths and finally collected on a bobbin by means of a Leesona fiber winder. The fibers are placed in water extraction baths for 24 h to remove residual solvent and finally air-dried under ambient conditions. The design of this spin line allows for the ability to monitor and control many of the spinning parameters that could potentially affect the electrical and mechanical properties of the polyaniline fibers. Furthermore, this design allows us to fabricate small lengths of fiber samples (<10 m) or prolonged runs in which tens of thousands of meters of the polyaniline monofilament is produced over a 24 h period. A photograph of the fiber production facility at SFST is shown in Figure 2.11. 2.3.3.2 Effect of Processing Conditions on the Physical Properties of the As-Spun Polyaniline Fibers Since there are many process steps involved in the formation of wet-spun fibers, the effect of these process parameters must be thoroughly investigated to determine the optimal set of conditions to maximize the desired mechanical properties of the as-spun polyaniline fibers. In particular, for spinning
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FIGURE 2.11 Photograph of the fiber production facility at SFST looking from the bobbin on the Leesona fiber winder toward the coagulation bath.
of polyaniline fibers, it is highly desirable to minimize or even prevent, the formation of macrovoids as they lower the mechanical strength. The formation of macrovoids is generally caused by fast precipitation kinetics of the fiber-spinning solution in the coagulation bath. To understand the macrovoid formation process when concentrated EB solutions were spun into a water coagulation bath, the effects of EB concentration, coagulation bath composition and temperature, and gel inhibitor additives, on the morphology of the resulting fibers were investigated [41]. In these studies, the as-spun polyaniline fibers were processed using water as the coagulant because it is cheap, safe, and easily treated after it is contaminated. The PANI fibers were initially processed according to the methods described previously using a take-up speed of 10 m=min on the first godet from different concentrations of EB dissolved in the NMP=HPMI mixture. Figure 2.12 shows the cross sections and the microporosity from one of the representative as-spun fibers collected from different fiber-spin trials. Macrovoids are observed for every fiber sample regardless of dope solution concentration. This result is consistent with previous studies in which the polyaniline fibers were spun from concentrated EB solutions dissolved in NMP [39] and DMPU [36]. Due to the formation of macrovoids in the as-spun fibers, the diameter of the fiber (250–300 mm) was always larger
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17.5 wt%
20 wt%
22 wt%
25 wt%
FIGURE 2.12 Effect of EB concentration on the morphology of cross section and microstructure of the resulting fibers spun into 258C water bath. (Reprinted from Yang, D., Fadeev, A.G., Adams, P.N., and Mattes, B.R., Proc. SPIE, 4329, 59, 2001. With permission from the International Society of Optical Engineering.)
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40
0.6
35 0.55 30 Macrovoids (%)
20
0.45
15
0.4
Fiber density (g/cm3)
0.5 25
10 0.35 5 0
0.3 16
18
20
22
24
26
EB concentration (wt%)
FIGURE 2.13 Effect of EB concentrations on the amount of macrovoids in the fiber cross section, the fiber density and the average length of macrovoids in the cross section. (Reprinted from Yang, D., Fadeev, A.G., Adams, P.N., and Mattes, B.R., Proc. SPIE, 4329, 59, 2002. With permission from the International Society of Optical Engineering.)
than that of the 100 mm diameter spinneret used to process the fiber. Figure 2.13 shows the effects of EB concentration on the percentage of macrovoids and the density of the as-spun fiber. As noted, the excellent miscibility between water and NMP may lead to an instantaneous demixing, which is responsible for macrovoid formation. However, as the polymer concentration increases, the viscosity of the EB solution increases. The high viscosity of the solution that easily gels reduces the solidification time of the dope solution. Hence, it is more difficult for water to diffuse into the center region of the nascent fibers. Therefore, the opportunity of macrovoid formation over a large area is limited. This is the reason why the amount of macrovoids in the fiber cross section decreases from 30% to 9% as the EB concentration increases from 17.5 to 25 wt%. Furthermore, the size of the macrovoids on the cross section decreases significantly with the increasing EB concentration from 50 to 20 mm. Although the diffusion of the water was slowed in the more concentrated EB solutions, water molecules, which have a very small size, can still penetrate into the center of the nascent fiber. However, the slow diffusion of water into the fiber reduced the amount of water diffusing into the central region of the as-spun fiber, which caused a more uniform porous structure rather than generating a large number of macrovoids. Therefore, the standard deviation of the pore size decreases from 0.65 to 0.35 mm with the EB concentration. As the number of macrovoids decreased, the density of the fibers increased. The stress–strain curves of these fibers are summarized in Figure 2.14. As expected, the fewer macrovoids the fibers possess, the stronger the fibers become. The fiber spun from the 17.5 wt% EB solution was too weak to be tested, and therefore, there are no mechanical properties reported for this batch of fibers. As the solid content of the dope solution increased from 20 to 25 wt%, the Young’s modulus increased 1.7 times whereas the tensile strength of the fibers increased from 5.9 to 12.8 MPa. However, all of these fibers are stiff and brittle, with a strain at break of only 1%. Consequently, they are too weak to be used for practical applications. From the morphology and mechanical properties of these as-spun PANI fibers, it is obvious that the size and number of macrovoids must be reduced. Therefore, the effects of the temperature and composition of the coagulation bath were subsequently investigated. The coagulation bath temperature should affect the morphology of the as-spun PANI fibers. From theoretical calculations, a high bath temperature is expected to accelerate the diffusion and the demixing processes. From the above study, we
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15
12
Stress (MPa)
20 wt% 22 wt% 25 wt% 9
6
Properties
3
20 wt% EB
22 wt% EB
25 wt% EB
Fiber diameter (µm)
246 ± 20
244 ± 7
249 ± 5
Modulus (MPa)
550 ± 29
684 ± 32
939 ± 2
Tensile strength (MPa)
5.9 ± 24
8.3 ± 4.2
12.8 ± 3.8
Elongation (%)
1.1 ± 0.5
6.4 ± 1.1
1.4 ± 0.6
0 0
0.5
1
1.5
2
Strain (%)
FIGURE 2.14 Effect of EB concentration on the mechanical properties of the fibers spun into the water bath at 258C. (Reprinted from Yang, D., Fadeev, A.G., Adams, P.N., and Mattes, B.R., Proc. SPIE, 4329, 59, 2003. With permission from the International Society of Optical Engineering.)
know that an instantaneous demixing occurs even at room temperature. At high temperature, the fast diffusivity of water allows the water molecules to penetrate deeper into the fibers and therefore, promotes macrovoid formation in the central region in addition to the edge. The EB fibers fabricated in this study were spun using a take-up speed of 15 m=min, and the temperature of the coagulation bath was set to 258C, 358C, and 458C. The temperature of the two godet baths was set at the same temperature as that of the coagulation bath. The cross sections and the microstructures of the obtained fibers (Figure 2.15) show that the water bath temperature significantly affects the morphology of the spun fibers. As expected, as the temperature increases from 258C to 458C, the amount of macrovoids in the fiber cross section increases from 9% to 17%, the average size of the macrovoids increases from 12 to 32 mm, the microporous size increases from 0.7 to 1.1 mm, and thus the diameter of these fibers increases from 153 to 240 mm. The mechanical properties of the PANI fibers spun using different coagulation bath temperatures are shown in Figure 2.16. At a coagulation bath temperature of 458C, the fiber possessed the lowest values for both tensile strength and Young’s modulus for all of the fibers spun as part of this study. This again suggests that the amount of macrovoids in the fiber influences their mechanical properties. Therefore, lowering the temperature of coagulation bath can minimize the formation of macrovoids for EB=NMP=HPMI solutions. Changing the coagulant can affect both the diffusion and the demixing processes between the solvent and the nonsolvent. A slow diffusion of the solvent into nonsolvent can reduce the probability of macrovoids forming in the central region of nascent fibers. The formation of macrovoids can be avoided by delaying demixing except when the delaying time is very short [47]. Adding solvent into the coagulation bath is a common technique to minimize macrovoid formation because the presence of solvent in the coagulation bath not only slows down the diffusion rate of the nonsolvent into solvent but also delays the demixing between solvent and nonsolvent. Using a 25 wt% EB=NMP=HPMI solution, we also spun a batch of fiber into a 20 wt% NMP aqueous bath. The PANI fibers were processed using coagulation bath temperature of 208C with a take-up speed on the first godet of 10 m=min. Typically, in the water bath, the nascent fiber formed within a few minutes. In order to remove any residual solvent (NMP and HPMI) from the as-spun fiber, we usually kept the fibers in a water bath for 24 h. However, for the same dope solution, which was spun into the
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25˚C
35˚C
45˚C
FIGURE 2.15 Effect of temperature on the morphology of the EB fibers spun from the 25 wt% EB=NMP=HPMI solution coagulated in the water bath. (Reprinted from Yang, D., Fadeev, A.G., Adams, P.N., and Mattes, B.R., Proc. SPIE, 4329, 59, 2004. With permission from the International Society of Optical Engineering.)
20 wt% NMP aqueous bath, it was impossible to collect the nascent fibers due to the severe deformation when the immersion time was shorter than a few minutes. This suggests that the precipitation rate of the dope solution in the 20 wt% NMP aqueous bath was significantly slowed due to the presence of NMP. However, since there was still a significant amount of water (80 wt%) in the coagulation bath, its presence dominates the formation of macrovoids. Therefore, the amount of macrovoids on the cross section for both fibers are similar to one another. Nevertheless, when the EB solution was coagulated
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15 25°c 35°c 45°c
Stress (MPa)
12
9 Properties
6
3
25˚C
35˚C
45˚C
Fiber density (g/cm3)
0.814 ± 0.011
0.560 ± 0.012
0.492 ± 0.008
Fiber diameter (µm)
153 ± 15
183 ± 0
240 ± 18
Modulus (MPa)
1245 ± 7.8
969 ± 113
788 ± 40
Tensile strength (MPa)
14.1 ± 7.8
9.4 ± 4.6
7.1 ± 2.6
1.1 ± 0.6
1.0 ± 0.4
0.9 ± 0.3
Elongation (%)
0 0
0.5
1
1.5
2
Strain (%)
FIGURE 2.16 Effect of temperature on the mechanical properties of the EB fibers collected from the water bath. (Reprinted from Yang, D., Fadeev, A.G., Adams, P.N., and Mattes, B.R., Proc. SPIE, 4329, 59, 2005. With permission from the International Society of Optical Engineering.)
into the 20% NMP aqueous bath, the solvent-exchange rate between NMP and water containing 20% NMP is slower than that between NMP and pure water. Thus, the penetration rate of water in the nascent fiber in the former situation is slower than that in the latter situation. Figure 2.17 shows the microstructure of the fiber made from the 20% NMP aqueous bath. It is more uniform and considerably denser than that of the fiber collected from the water bath. Comparing the fiber collected from 20% NMP aqueous bath to the fiber collected from the water bath, Figure 2.18 shows that the former has a smaller diameter and a higher density and thus possesses a higher tensile strength and a tougher strain at break than the latter. The PANI fiber spun into the 20 wt% NMP aqueous coagulation bath possessed a Young’s modulus of 1.3 GPa and a tensile strength of 30 MPa whereas the PANI fiber spun into the water coagulation bath possessed a Young’s modulus of 940 MPa and a tensile strength of 23 MPa. For comparison with PANI fibers spun from concentrated EB=NMP=HPMI solutions, we also carried out an analogous study on fiber spinning from an EB=NMP=4MP solution. The effect of the coagulant temperature was included in this study because 4MP is more water soluble than HPMI. Therefore, the temperature of the aqueous coagulation bath may have a more pronounced impact on the fiber morphology. PANI fibers were spun from a 20 wt% EB=NMP=4MP solution with temperatures of 208C or 358C for the water coagulation bath and godet baths. The take-up speed of the PANI fiber on the first godet was 10 m=min. The morphology of two fiber samples spun at the different coagulation bath temperatures is shown in Figure 2.19. Interestingly, it was found that the effect of the coagulation bath temperature on the morphology of the fibers was opposite to that of the fibers spun from concentrated EB=NMP=HPMI solutions in that fewer macrovoids were formed when the coagulant temperature was raised. Figure 2.19 indicates that a fiber nearly free from macrovoids was formed at higher coagulation and godet bath temperatures. It is apparent that the coagulation bath temperature has a different effect for the EB=NMP=4MP dope solution than for the EB=NMP=HPMI dope solution. It is postulated that 4MP, a much smaller molecule, can diffuse more rapidly from the polymer solution than HPMI. This promotes fiber solidification and increased coagulation rate compared to the EB=NMP=HPMI solution. When the coagulant bath temperature was elevated, the diffusion rate of 4MP and its water solubility increased. Consequently, the dope solution precipitated more rapidly, leading to nearly macrovoid-free fibers.
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Water bath
20wt% NMP aqueous bath
FIGURE 2.17 Cross sections and micropore structures of the fiber samples in water and the 20 wt% NMP aqueous coagulation baths. (Reprinted from Yang, D., Fadeev, A.G., Adams, P.N., and Mattes, B.R., Proc. SPIE, 4329, 59, 2006. With permission from the International Society of Optical Engineering.)
The mechanical test results for fiber spun using a coagulant bath of EB=NMP=4MP at temperatures between 208C and 358C are shown in Figure 2.20. The EB fibers spun using a coagulant temperature of 358C had a modulus of 3.5 GPa, a tensile strength of 44 MPa, which is stronger than the EB fiber spun using a coagulant temperature of 208C (modulus of 1.2 GPa, tensile strength of 28 MPa). However, the strain of the latter fiber was four times greater. It should be noted that the tensile strength and modulus of EB fibers spun using a coagulant temperature of 358C are significantly higher than for unstretched EB fibers reported by other researchers [33,39]. However, in order to make practical devices from the base-processed fibers, these mechanical properties still need to be improved.
2.3.4 Asymmetric Polyaniline Hollow Fibers for Membrane Applications Commercial scale processing of polyaniline into integrally skinned asymmetric hollow fiber membranes is more technically challenging than for other polymeric hollow fibers due to the inherent difficulties of processing concentrated polyaniline solutions for the reasons outlined above for producing polyaniline textile fibers. However, fabrication of asymmetric hollow fiber membranes from concentrated, highmolecular-weight emeraldine base solutions can be achieved through the use of small amounts of secondary amine additives to stabilize the concentrated polyaniline solutions for over 30 h. This increased processing window enables the commercial scale processing of polyaniline into hollow fiber
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35 20% NMP/Water
30
Water bath
Stress (MPa)
25 Properties
20
Water bath
Macrovoid (%)
6
7
Length of macrovoids (µm) 16.2 ± 7.8
15
Pore size (µm) Fiber density (g/cm3)
10
0
1
2
1.55 ± 0.61
1.31 ± 0.41
0.574 ± 0.006
0.738 ± 0.021 189 ± 3
Modulus (MPa)
943 ± 46
1331 ± 216
Tensile strength (MPa)
23.4 ± 2.9
29.6 ± 2.9
4.0 ± 1.1
6.4 ± 1.1
Elongation (%)
0
14.6 ± 5.9
267 ± 2
Fiber diameter (µm) 5
20wt% NMP aqueous bath
3 Strain (%)
4
5
6
FIGURE 2.18 Stress–strain curves for the EB fiber precipitated in water and the 20 wt% NMP aqueous coagulation baths. (Reprinted from Yang, D., Fadeev, A.G., Adams, P.N., and Mattes, B.R., Proc. SPIE, 4329, 59, 2007. With permission from the International Society of Optical Engineering.)
20˚C
FIGURE 2.19
35˚C
Morphology of the EB fibers precipitated in a water coagulation bath at different temperatures.
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50 Fiber spun at 20˚C
Stress (MPa)
40
Fiber spun at 35˚C
30
20
10
0 0
2
6 4 Strain (%)
8
10
FIGURE 2.20 Stress–strain curves of the EB fibers precipitated in a water coagulation bath at different temperatures.
membranes. Using this emeraldine base-processing route, we have shown in our laboratory that it is possible to control the phase inversion process in order to fabricate polyaniline hollow fiber membranes with different mesoscopic morphologies. The mesoscopic morphology of these hollow fibers defines both their mechanical properties and usefulness for molecular discrimination. Interest in polyaniline membranes for gas separation stems mostly from the work by Anderson et al. [48,49], which reported some of the highest ideal separation factors obtained with polymeric membranes for O2=N2 (30), H2=N2 (3590), and CO2=CH4 (336). They additionally showed that the selectivity and permeability could be altered by doping or dedoping and redoping the polyaniline membranes with different acids. However, despite these high separation factors, their commercial viability is still questionable due to the orders of magnitude lower permeability of dense polyaniline membranes compared to other polymeric membranes. One of the most widely used approaches to increase the flux through a membrane is to deposit the ultrathin-selective film onto a highly permeable support membrane (i.e., a thin-film composite membrane). This approach was first reported by Kuwabata and Martin [50], and utilized thin polyaniline films deposited onto alumina support membranes and showed a large increase in the total gas permeation rate. It is well known that the flux is inversely proportional to the film thickness. Subsequently, Lee et al. [51] used a porous nylon support membrane for improving the permeation of N2 and O2 through a partially redoped polyaniline membrane and obtained an ideal separation factor of 28. While these previous studies have used polyaniline flat-sheet membranes, hollow fiber membranes offer the advantage of manufacturing into membrane modules with a 4–10 times increase in membrane surface area per unit volume. It has been previously demonstrated that polyaniline hollow fibers can be made with a uniformly dense structure on a laboratory scale using an acid-processing route [52]. However, it is preferred that the hollow fiber is formed as a microporous structure with a thin, dense selective layer to increase the gas permeance through the membrane. In our laboratory, asymmetric polyaniline hollow fiber membranes of varying morphologies for gas separations have been fabricated. The outer diameter of these hollow fibers was chosen to be less than 500 mm since hollow fibers with these dimensions are typically used for low and medium pressure applications (feed pressure: 50–100 psi). Norris et al. [53] showed that by adjusting the processing parameters it is possible to control the thickness, porosity, and location of the ‘‘skin layer’’ to provide a
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range of morphologies from a highly permeable microporous interface to a dense ‘‘barrier layer.’’ The thickness of the skin layer was controllably adjusted between 0.5 and 5 mm, and hence enables these asymmetric integrally skinned polyaniline hollow fiber membrane to be used for gas separations. The skin layer can be formed on either the outer surface of the hollow fiber or in the inner lumen depending on the final application of the hollow fiber membrane. 2.3.4.1 Spinning of Asymmetric Hollow Fiber Membranes To successfully spin polyaniline into asymmetric hollow fibers, the minimum solids content required to prevent the hollow fiber from becoming deformed or collapsing in the coagulation bath was found to be 15 wt%. The polyaniline hollow fibers were processed from emeraldine base solutions in which 18 wt% of high-molecular-weight polyaniline emeraldine base powder (Mw 300,000 g=mol) was dissolved in NMP containing the gel inhibitor, 4-methylpiperidine. The stoichiometric ratio between gel inhibitor molecules and the EB tetramer repeat unit required to stabilize the dope solution was 1.2:1. This produced a fiber-spinning dope solution that possessed a stable viscosity of approximately 150,000 cP for at least 6 h. The gel inhibitor 4-methylpiperidine results in only a modest reduction in the oxidation state in the polymer, whereas more aggressive secondary amine gel inhibitors, such as 2-methylaziridine, were found to reduce the polymer at substantially faster rates and hence were unsuitable for hollow fiber production on a continuous basis. The concentrated polyaniline solution was extruded through the outer annulus (530 mm OD) of the hollow fiber spinneret, whereas the bore fluid was delivered through the inner capillary (200 mm OD). Depending on whether the skin layer was formed on the outside surface or on the inner lumen, the polymer solution was either extruded through an air gap or steam zone into the coagulation bath. The take-up speed on the first godet was set to between 120 and 150 m=h. The nascent fiber passed through two washing godet baths and finally collected on a bobbin by means of a Leesona fiber winder. The compositions of the bore fluid, coagulation bath, and the godet baths can be varied in order to obtain different membrane morphologies. Generally, the coagulation and godet baths were filled with water or acidic solutions with a pH less than 3. Acidic coagulants or washing solutions are used so that the EB hollow fiber became fully doped with the desired dopant acid by the time the hollow fiber is collected onto the bobbin. This step eliminates any subsequent postprocessing step to dope the fiber with the desired acid. In contrast, the bore fluid can consist of water and organic solvent mixtures (isopropanol–water, ethanol–water, acetone–water) or acidic solutions (aqueous phosphoric acid solutions). The reason for these different bore fluid compositions is to control the location and thickness of the dense, anisotropic skin as they influence the precipitation kinetics of the polymer solution. The polyaniline hollow fiber membranes collected on the bobbin are generally immersed in methanol extraction baths to remove any residual solvent and the gel inhibitor before drying under ambient conditions. Methanol was chosen instead of water for this extraction process because water could potentially collapse the pores upon drying due to capillary forces. 2.3.4.2 Properties of Asymmetric Hollow Fiber Membranes For gas separation applications, the feed stream is usually fed into the shell side of the module. This implies that the dense selective layer should be located on the outside of the polyaniline hollow fiber membrane. To achieve this desired morphology, the polyaniline hollow fiber was spun using an air gap between the spinneret and the coagulation bath. The residence time in the air gap influences the amount of solvent evaporation, which in turn governs the thickness of the dense separating layer on the outer surface of the hollow fiber. By adjusting the residence time from a few seconds to 30 s, the thickness of the dense separating layer on the outer surface of the hollow was successfully varied between 0.5 and 5 mm. The formation of macrovoids in these asymmetric polyaniline hollow fibers is undesirable as they weaken the mechanical strength and may lead to defects in the selective layer, which make them nonviable for many separations applications. The formation of macrovoids in hollow fiber membranes is generally caused by fast precipitation kinetics of the polymer solution in the coagulation bath. The formation of macrovoids when spinning polyaniline hollow fibers was found to be highly dependent on
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the water content in the bore fluid. As the water content of the bore fluid decreased, the number and size of the macrovoids in the hollow fibers decreased. Bore fluids that consisted of isopropanol–water mixtures were found to be superior to both ethanol–water and acetone–water mixtures for producing asymmetric hollow fibers with the least number of macrovoids at the same organic solvent concentration. The hollow fiber shown in Figure 2.21 was spun using a 40 wt% isopropanol–water solution as the bore fluid into a water coagulation bath. Using a residence time of 2 s in the air gap between the spinneret and the coagulation bath, we were able to fabricate a hollow fiber with a 500 nm thick dense skin on the outer surface and a microporous inner surface. This hollow fiber had an outer diameter of 480 mm and an inner diameter of 280 mm. Although a small amount of macrovoids was formed under the dense skin, it is expected that this should not sacrifice the gas selectivity of this membrane. Furthermore, it is expected that this hollow fiber will possess a higher gas permeance than previously reported polyaniline hollow fibers prepared from an acid-processing route by Monkman and coworkers [52], which had dense walls with a thickness of 90 mm (H2 permeability of 4 Barrers). It has been previously shown that the type of acid used to dope polyaniline flat sheet membranes affects both their permeability and selectivity [48]. The selective nature of these hollow fiber membranes can be tailored during the fiber-spinning process by doping the polyaniline hollow fiber with the desired acid. The godet baths were filled with a 1 M acidic solution of varying acid strength and size, and the
(a)
(b)
(c)
FIGURE 2.21 Morphology of an asymmetric polyaniline hollow fiber membrane with an 800 nm dense skin on the outer surface: (a) overall cross section of the hollow fiber (magnification ¼ 180), (b) cross section of the outer surface (magnification ¼ 10,000), and (c) cross section of the fiber wall (magnification ¼ 10,000). (Reprinted from Norris, I.D., Fadeev, A.G., Pellegrino, J., and Mattes, B.R., Synth. Met., 153, 57, 2005. With permission from Elsevier.)
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Conjugated Polymers: Processing and Applications TABLE 2.5 Resistance of the Base-Processed Polyaniline Hollow Fibers Doped with Different Acids during the Fiber-Spinning Process Doping Acid
pKa
Phosphoric acid HBF4 HCl Pyruvic acid Acrylic acid
2.13 0.5 7 2.35 4.29
Fiber Resistance (V=cm Length of Fiber) 140 130 140 160 130
total residence time in each of the acidic godet baths was 10 min. The effectiveness of the doping was determined by measuring the resistance of the fiber after drying under ambient conditions for at least 24 h. Table 2.5 shows that the resistances per unit length of these hollow fibers were essentially identical within experimental error, which implies that this residence time in the godet baths was sufficient to fully dope the polyaniline hollow fibers.
2.4 2.4.1
Acid-Processing Route for the Fabrication of Polyaniline Fibers Introduction
Besides using a base-processing route for polyaniline fiber preparation, it is also possible to process polyaniline fiber in its electrically conductive, emeraldine salt form. Base-processed fibers require a postprocessing acid-doping treatment to make the polyaniline fibers electrically conductive. The acidprocessing route has several advantages including dopant homogeneity in the final product, as compared to the inhomogenous doping typically encountered in base-processed polyaniline fibers. Additionally, the mechanical properties of base-processed fibers are often adversely affected by the doping process. The earliest example of using an acid-processing route was reported by Andreatta and coworkers [54,55] in which polyaniline fibers were prepared from a 20 wt% solution of polyaniline dissolved in 96% sulfuric acid. The concentrated polyaniline solution was extruded at 608C, and dry-jet wet-spun into chilled water. The electrical conductivity of the as-spun polyaniline fibers was determined to be in the range from 20 to 60 S=cm and fibers displayed sharp x-ray reflections, which indicates a significant degree of crystallinity. Cao et al. [56] showed that processing of polyaniline in the conducting emeraldine salt form can be achieved by doping emeraldine base with an acidic counterion with ‘‘surfactant-like’’ properties. Camphorsulfonic acid (CSA) and dodecylbenzenesulfonic acid have proven to be successful ‘‘surfactant-like’’ dopants to render polyaniline soluble in a wide range of organic solvents, including NMP, chloroform, xylene, formic acid, dimethyl sulfoxide, dimethyl formamide, and m-cresol. It was surmised that the anionic part of the acid dopes the polyaniline whereas the long alkyl chain leads to solubility in the organic solvent. In particular, solutions of camphorsulfonic acid-doped polyaniline, when cast into free-standing films from m-cresol, are of considerable interest due to the combined effect of high conductivity (100–400 S=cm) and good mechanical properties. However, m-cresol is a toxic solvent, and considerably difficult to remove from the polyaniline films (or fibers). Building upon this concept, Wang et al. [57] showed that conductive fibers could be prepared from concentrated solutions of polyaniline and poly(o-toluidine) doped with CSA in m-cresol. The as-spun polyaniline fiber was prepared from a 28.0 wt% mixture of polyaniline doped with CSA dissolved in m-cresol. The solution was extruded through a 500 mm diameter single-hole spinneret into an ethyl acetate coagulation bath. The tensile strength, elongation, and modulus of the as-spun polyaniline fiber were found to be 0.2 gpd, 8.4%, and 7.3 gpd, respectively. Similarly, the poly(o-toluidine) fiber was prepared from a 30.2 wt% mixture of poly(o-toluidine) doped with CSA dissolved in m-cresol (a highly carcinogenic solvent). The solution was extruded through a 500 mm single-hole spinneret into a toluene
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coagulation bath. The tensile strength, elongation, and modulus of the as-spun poly(o-toluidine) fiber were found to be 0.2 gpd, 3.0%, and 9.7 gpd, respectively. The conductivity for the as-spun polyaniline fiber was found to be 200 S=cm. With the substitution of a larger methyl group for a hydrogen atom in poly(o-toluidine), more disorder is introduced into the as-spun poly(o-toluidine) fiber and this reduced its conductivity to 10 S=cm. This study also measured the transport properties (temperature-dependent conductivity, thermopower, microwave frequency conductivity, and dielectric constant) for the as-spun polyaniline fibers and found that the as-spun polyaniline fiber showed similar charge transport behavior to camphorsulfonic acid-doped polyaniline films cast from m-cresol, although the fiber did contain some areas of poorly conductive linkages. It should be noted that the highly toxic nature of m-cresol prevents commercial processing of polyaniline fibers using this solvent. In the work by Cao et al. [56], it was observed that the films of polyaniline doped with camphorsulfonic acid, when cast from solutions of the polymer dissolved in dichloroacetic acid (DCAA) possessed a conductivity of 80 S=cm. Although this film did not possess the same level of conductivity as the films processed from phenol solvents, it was higher than for films cast from other solvents including NMP, chloroform, and dimethyl sulfoxide. Based on this observation, Monkman and coworkers [58] developed a new acid solution-processing route for preparing polyaniline films via a new route comprised of 2-acrylamido-2-methyl-1-propane sulfonic acid (AMPSA) as both the protonating acid and the solvating group, and DCAA as the solvent. It was possible to draw these AMPSA-doped films uniaxially both at room temperature and at 908C. The room temperature conductivity along the stretch direction was increased to a maximum value of 670 S=cm for a film drawn at 908C, compared to 210 S=cm for the ascast film. More recently, Monkman and coworkers [59–61] showed that this acid solution-processing route for polyaniline also enables the wet spinning of electrically conductive polyaniline fibers. The fibers were processed from 9 wt% solutions of high-molecular-weight emeraldine base (Mw 150,000 g=mol) protonated with AMPSA dissolved in DCAA. The ratio of the number of AMPSA molecules to the number of nitrogen atoms in the polyaniline was typically 0.6 (2.4 moles per EB tetramer repeat unit) since this molar ratio was shown to produce polyaniline films from these solutions with the highest electrical conductivity. The fibers were spun into various coagulation solvents. The three preferred coagulants were acetone, butyl acetate, and 4-methyl-2-pentanone. The as-spun polyaniline fibers have a modulus of 40–60 MPa, tensile strengths of 20–60 MPa, and electrical conductivities of 70–130 S=cm, which increased to 1000 S=cm after stretching the as-spun fiber over a 908C hot pin to five times its initial length. It was also shown that drawing the as-spun fibers at elevated temperatures also enhanced the mechanical properties of the fibers. For example, the polyaniline fibers drawn at 908C possessed a modulus of 2 GPa and a tensile strength 97 MPa, while retaining conductivities of 600 S=cm. Spinks et al. [62] later modified this process to form composite polyaniline fibers to fabricate highstrength ICP-based artificial muscles. The composite fibers were spun from a concentrated solution of AMPSA-doped polyaniline dissolved in DCAA that also fully dispersed single-walled carbon nanotubes. The role of the carbon nanotubes was to increase both the mechanical and electrical properties of the AMPSA-doped polyaniline fiber since the addition of carbon nanotubes to other polymer matrices has yielded materials with increased stiffness and strength. The spinning solution was extruded through a 250 mm diameter spinneret into an acetone coagulation bath. The spun fibers were stretched to approximately five times their original length across a soldering iron wrapped in Teflon tape heated to 1008C. The addition of 0.76 wt% carbon nanotubes to the AMPSA-doped polyaniline fiber resulted in the tensile strength of the polyaniline fiber increasing from 170 to 255 MPa and the modulus increasing from 3.4 to 7.3 GPa. Similarly, the conductivity of the polyaniline fiber increased from 497 to 716 S=cm.
2.4.2 Recent Advances in the Production of Acid-Processed Polyaniline Textile Fibers The mechanical and electrical properties for the polyaniline fibers spun using the methods published by Monkman and coworkers [59–61] are encouraging as they approach those of conventional textile fibers such as nylon-6. However, in our laboratory it was found that direct implementation of the reported
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procedure for preparing concentrated solutions of AMPSA-doped polyaniline dissolved in DCAA (PANI–AMPSA=DCAA) was found to be unsuitable for preparing fibers on a pilot plant scale due to rapid gelation of the dope solution. It was noted by Monkman and coworkers [61] that the maximum solids content (EB and AMPSA) at which gelation is not experienced is 5 wt%, and these solutions did not form a continuous fiber. The maximum solids content, which they could fabricate before the solution became too viscous to enable fiber spinning due to rapid gelation of the dope solution was 10 wt%. Commercial scale processing of textile fibers generally requires that the solids content in the fiber-spinning dope solution be above 10 wt%. Therefore, it was necessary to explore the preparation of stable dope solutions with greater than 10 wt% total solids that would allow continuous fiber production for at least 24 h. By precisely controlling the water concentration, unlike the anhydrous preparation procedure developed by Monkman and coworkers [59–61], it was found that up to 14 wt% PANI– AMPSA=DCAA dope solutions could be prepared, which were stable for several months when stored below 58C. Conducting polyaniline fibers with a modulus of 3–6 GPa, peak stress of 120–200 MPa, extension at break of 10%–20%, and conductivity up to 900 S=cm have been fabricated from these 12 to 14 wt% PANI–AMPSA=DCAA solutions on a pilot plant scale [63]. Based on this research, SFST has introduced these highly conductive polyaniline fibers to the marketplace under the trade name of Panion. This combination of mechanical properties makes these Panion monofilaments suitable for weaving, knitting, stitching, and braiding using conventional textile-processing equipment. For example, Figure 2.22 shows scanning electron micrographs (SEM) of Panion monofilaments that are processed into yarns as well as knitted and three-dimensional braided textile structures using conventionalprocessing equipment. 2.4.2.1 Preparation of Polyaniline Dope Solutions It is highly desirable to raise both the solid content in the dope solutions and the molecular weight of the polyaniline from which the solutions are prepared, as this raise should produce fibers with better mechanical and electrical properties than those fabricated by Monkman and coworkers [59–61]. Attempts to prepare concentrated polyaniline dope solutions with the high-molecular-weight polyaniline synthesized in-house (Mw 280,000 g=mol) using the mixing protocol described by Monkman and coworkers [59–61] were unsuccessful due to the formation of a thermoirreversible gel once the total solids (EB and AMPSA) concentration exceeds 4.5 wt%. It appears that the cause of this gelation is related to the mixing process, which generates a significant amount of heat that causes the cross-linking of the AMPSA molecules through the C¼C bond. For example, the temperature after the 4.5 wt% PANI– AMPSA=DCAA solution was homogenized for 30 min was 708C. This suggests that the temperature of the dope solution requires control during the mixing process. The formation of a thermoirreversible gel has to be overcome in order to spin polyaniline fibers on a pilot plant scale for at least 24 h using this high-molecular-weight, acid-processing route. Monitoring of the solution temperature during mixing showed that gelation occurred if the solution was heated above 588C. Maintaining the mixing temperature below 408C ensured that the solutions did not form a thermoirreversible gel and remained fluid during spinning. It was found that maintaining the temperature of a 4.5 wt% PANI–AMPSA=DCAA solution below 408C using a 108C water bath resulted in a threefold reduction in the room temperature viscosity (12,000–4000 cP) and extended the lifetime of the solution before gelation from less than a few hours to more than several days. In the preparation of these concentrated PANI–AMPSA=DCAA solutions according to the procedures reported by Monkman and coworkers [59–61], it was noticed that the high-molecular-weight EB powders, which had been dried in the vacuum oven (water content <0.5 wt%), had difficulty wetting upon addition to the DCAA solvent. This problem could be overcome by the addition of a trace amount of water to the EB powder before it was dissolved in the DCAA solvent. However, there was a finite water limit. Above 30 wt% water in EB powder, the nascent fiber flattened and fell apart in the coagulation bath. Furthermore, it was found that fiber with good mechanical properties could only be spun if the water content in the EB powders was between 2 and 10 wt%. Thus, with appropriate mixing conditions, PANI– AMPSA=DCAA dope solutions with a total solids content between 4.5 and 14 wt% may be prepared
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(a)
(b)
(c) FIGURE 2.22 SEM micrographs of (a) yarn formed from 20 Panion monofilaments having a twist ratio of 7 turns per inch, (b) 40 denier Panion monofilaments knitted into a fabric with woven nylon fabric, and (c) tubular braid of 40 denier Panion monofilaments.
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without the solution undergoing gelation with a 24 h period. These solutions were prepared using the high-molecular-weight polyaniline 8,000 synthesized in-house (Mw 280,000 g=mol) from phosphoric acid as described in Section 2.2.3. The ratio of the number of AMPSA 6,000 molecules to the number of nitrogen atoms in the polyaniline was 0.6 (2.4 moles per EB 4,000 tetramer repeat unit). As with the base-processed fiber, understanding the rheology of the fiber-spinning 2,000 dope solutions is important to enable the fabrication of polyaniline fibers with repro0 ducible electrical and mechanical properties. 0 3 6 9 12 15 18 Figure 2.23 illustrates the viscosity behavior Time (h) of a 7 wt% PANI–AMPSA=DCAA dope soluFIGURE 2.23 Rheological behavior of a 7 wt% total solids tion over an 18 h period when measured PANI–AMPSA=DCAA solution. using a Brookfield RVDV III cone and plate. The viscosity of this solution was measured at 258C and a constant shear rate of 0.8 s1. Three distinct rheological regions were observed for these polyaniline dope solutions. Region I corresponds to the dissolution of EB into the DCAA solvent containing the AMPSA dopant with the viscosity of the mixture decreasing with the shearing time. Region II is an equilibrium state. As more polyaniline chains take on an expanded conformation, a three-dimensional network starts to form between PANI and AMPSA molecules. In this region, under the shear rate of 0.8 s1, the rate of breaking hydrogen bonds between PANI chains is comparable to the rate of forming the hydrogen bonds between AMPSA and the PANI chains. It should be noted that the AMPSA may interact via ionic bonding between its sulfonic acid group and the protonated amine nitrogens. Alternately, the polar groups can also hydrogen bond with the unprotonated amine nitrogens in the polyaniline backbone. Therefore, the viscosity of the dope solution is fairly constant (3300 cP). The fiber spun from Region II solutions gives strong mechanical integrity and high conductivity. The length of Region II is highly dependent on the dope solution mixing condition. Although not shown in Figure 2.23, Region III is associated with the gelation of the polymer solution. It was found that once polyaniline dope solution showed signs of gelling, the spinning of the fiber with consistent mechanical properties becomes increasingly difficult. Although the fiber spun from a partially gelled dope solution possessed excellent modulus (3–6 GPa) and peak stress (100–200 MPa), the fiber is brittle (elongation at break <5%). Similar rheological studies of the spinning solutions indicate that the viscosity for 12 wt% total solids dope solutions is between 120,000 and 180,000 cP. When the mixing conditions are controlled, the rheological properties of the 12 wt% total solids dope solutions experience no significant changes when stored for 4 months at 58C and were found to be only partially gelled after 9 months of storage. Viscosity (cP)
10,000
2.4.2.2 Spinning of Doped Polyaniline Fibers The fibers that are described in the following section were spun from a 12 wt% PANI–AMPSA=DCAA dope solution with the high-molecular-weight polyaniline synthesized in-house (Mw 280,000 g=mol). The water content relative to the EB powder was between 2 and 10 wt% since these solutions yield fibers with good mechanical and electrical properties. This corresponds to an overall percentage of water in these PANI–AMPSA=DCAA solutions between 0.1 and 0.6 wt%. The ratio of the number of AMPSA molecules to the number of nitrogen atoms in the polyaniline was 0.6. A description of the spin line is identical to the one described earlier for fabricating the polyaniline fibers using the base-processing route (Section 2.3.3). In contrast to the base-processed polyaniline fibers described earlier, it was not possible to use water as a coagulant due to the slow precipitation kinetics of the dope solution. It has been generally found that
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esters, ketones, and alcohols in which DCAA is miscible, but is a nonsolvent for polyaniline, such as butyl acetate, acetone, methylisobutyl ketone, 2-butanone, methanol, ethanol, and isopropyl alcohol, can be employed to coagulate the dope solution [59–61]. However, ethyl acetate was found to be a superior coagulant to the organic solvents listed above in our pilot plant spin runs. It was observed that this coagulant yielded the best fiber in terms of both the mechanical and electrical properties of the asspun fiber and the fiber diameter was more uniform along the length of the fiber. The majority of studies [30,37–39,59–61] devoted to the stretching of the as-spun polyaniline fibers to increase their mechanical and electrical properties have been performed as independent postspinning processes rather than incorporated into the spinning process. In these studies, the as-spun polyaniline fiber is typically stretched across a heated soldering iron tip that was wrapped with a piece of Teflon film at temperatures between 908C and 2208C under slight uniaxial tension. As the heat softened the fiber, a draw stretch ratio of 3 to 5 times could be obtained. This method for aligning the polymer chains in the as-spun fibers is unpractical on a pilot plant scale and we modified the spin line in our laboratory to include a 2 m long heat tube maintained at a temperature between 908C and 1008C, which is located between the first and second godet baths. Stretching could be imparted on the fiber by rotating the stainless steel drums of the second godet faster than the first godet pair. Differential scanning calorimetry (DSC) characterization of the nascent fibers spun from concentrated PANI–AMPSA=DCAA dope solutions performed in our laboratory and by other researchers [61] shows an endothermic transition centered at 408C to 308C depending on the amount of residual DCAA in the fiber. The DSC trace above 58C is rather featureless with a gentle rise that is indicative of a plasticized system. This indicates that unlike camphorsulfonic acid-doped polyaniline films and fibers, which can only be drawn at temperatures 1508C [64], the fibers can be potentially stretched at room temperature. The optimal electrical and mechanical properties of the stretched fibers were obtained when the temperature inside the heat tube was 908C. It has been noted that the presence of AMPSA in the fiber is responsible for this plasticization, and above this temperature the AMPSA molecules that are hydrogen bonded to the polyaniline chains become increasingly mobile. This in turn enables undesirable motion of the polyaniline chains [61]. Two procedures for spinning the concentrated PANI–AMPSA=DCAA dope solutions into fibers were developed in our laboratory. In the first method, ethyl acetate was used as the coagulant. The dope solution was extruded through the spinneret (150 or 100 mm diameter single-hole spinneret) with an empty first godet bath. Consequently, the fiber that was collected on the bobbin was an AMPSA-doped polyaniline fiber. Since the performance of ICP-based devices is known to be dependent on the properties of dopant anions, for certain applications it may be desirable to replace the acid dopant ions, AMPSA and DCAA. Therefore, in our alternate method, ethyl acetate is still used as the coagulant, but the nascent fiber is immersed either in an acidic solution (e.g., phosphoric acid) or in basic solution (e.g., ammonium hydroxide) present in the first godet bath in order to modify the composition of the fiber by either exchanging (i.e., acidic solutions) or entirely removing (i.e., basic solutions) the AMPSA and DCAA dopants. In contrast to the polyaniline fibers that have been produced using the previously described baseprocessing route, one of the advantages of spinning polyaniline fibers using PANI–AMPSA=DCAA dope solutions is that the resulting fibers are dense and contain no macrovoids, which might reduce the mechanical properties of the fibers. Figure 2.24 shows an SEM image of a 30 mm AMPSA-doped polyaniline fiber produced at SFST. 2.4.2.3 Effect of Processing Conditions on the Physical Properties of the As-Spun Polyaniline Fibers In the work published by Monkman and coworkers [59–61], the polyaniline fibers were wet-spun into the coagulation solvent without applying any take-up mechanism or controlling the extrusion rate of the dope solution through the spinneret. This method for fiber spinning will inevitably result in some variation in the physical properties of the fibers due to varying extrusion rates and time spent in the coagulation bath. As noted earlier, there are many process steps involved in the formation of wet-spun
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FIGURE 2.24
SEM micrograph of a 30 mm AMPSA-doped polyaniline fiber spun at SFST.
polyaniline fibers and the effect of these process parameters must be thoroughly investigated to determine the optimal set of conditions that will maximize the desired mechanical and electrical properties of the as-spun polyaniline fibers. The polyaniline fibers were produced from a 12 wt% PANI–AMPSA=DCAA dope solution that was extruded through a 150 mm diameter spinneret into an ethyl acetate coagulation bath. The mechanical and electrical properties of these AMPSA-doped polyaniline fibers were found to be dependent on the extent of polymer chain alignment introduced into the fiber by adjusting the stretch ratio applied between the first and second godets. In these experiments, the stretch ratio between the first and second godet was increased from 1.2:1 to 2.7:1. The physical properties of the as-spun fibers with increasing amounts of stretch between the godets are shown in Table 2.6. It was found that continuous fiber spinning, under these conditions, required a stretch ratio 3:1. The amount of stretch that the fiber could withstand without breaking is at the lower end that was reported for AMPSA-doped polyaniline fibers that have been stretched over a heated soldering iron tip [59–61]. As expected, Table 2.6 shows that the fiber diameter and elongation decrease proportionally
TABLE 2.6
Variation in Properties of the As-Spun Polyaniline Fiber with Increasing Godet Stretch Ratio
Speed Ratio
Diameter (mm)
Density (g=cm3)
Denier (g=9000 m)
1.2:1 1.5:1 1.8:1 2.1:1 2.4:1 2.7:1
93 84 78 70 66 62
1.68 1.59 1.57 1.59 1.55 1.60
103 79 68 55 48 43
Conductivity (S=cm) 335 445 525 630 750 810
+ + + + + +
25 40 4 65 80 100
Tensile Strength (MPa) 62 77 70 79 97 111
+ + + + + +
5 5 8 2 5 3
Modulus (GPa) 0.6 1.1 1.4 2.1 2.6 2.9
+ + + + + +
0.1 0.1 0.1 0.1 0.2 0.4
Elongation (%) 84 60 42 29 14 12
+ + + + + +
8 4 10 10 4 4
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with increasing stretch ratio. This suggests that the fiber chains are increasingly aligned as the stretch ratio increases. As the stretch ratio increases, the denier (linear density of the fiber expressed in grams per 9000 m of fiber) also decreases, but the overall density of the fiber (averaging 1.60 + 0.08 g=cm3) does not change within experimental error (within 5%). As expected, Table 2.6 also shows that the tensile strength, modulus, and room temperature electrical conductivity increase with increasing stretch ratio, which again indicates increasing polymer chain alignment. The denier of the as-spun polyaniline fiber can be controlled not only by the take-up speed of the fiber on the first godet and the stretch ratio between the godets, but also by the diameter of the spinneret hole. A further set of experiments was carried out with a 100 mm diameter spinneret. These AMPSA-doped polyaniline fibers were produced from a 12 wt% dope solution that was extruded through a 100 mm diameter spinneret into an ethyl acetate coagulation bath. Besides determining the physical properties, the effect of the immersion time in the ethyl acetate coagulation bath was also investigated. These fibers were stretched 2.5 times between the godets, and the resulting properties are set forth in Table 2.7. The diameter for these as-spun fibers listed in Table 2.7 are between 43% and 45% of the spinneret diameter, which is similar to the ratio of those fibers highlighted in Table 2.6 stretched 2.4 and 2.7 times between the two godets having 44% and 41% of the spinneret diameter, respectively. In contrast to the properties of the fibers listed in Table 2.6, these fibers are observed to have slightly higher tensile strength and modulus whereas the percent elongation is lower. However, the densities of fibers spun with a residence time in the ethyl acetate coagulation bath of 115 and 36 s are slightly lower than with a residence time in the ethyl acetate coagulation bath of 17 s. The density of the latter fiber is similar to the values listed in Table 2.7. This reduced density may be due to the longer residence time in the coagulation bath, which allows ethyl acetate to penetrate into the nascent fiber to create microvoids, although this cannot be confirmed by SEM observation. It should be noted that polyaniline fiber spun with the residence time in the ethyl acetate coagulation bath of 17 s possessed the highest values for the modulus, tensile strength, and conductivity of the three samples listed in Table 2.7, likely due to higher density. Elemental analysis by energy dispersive x-ray spectroscopy (EDS) of the AMPSA-doped polyaniline fibers listed in Table 2.7 shows minimal extraction of AMPSA from the fiber during the spinning process. For EDS analysis, the fiber sample is subjected to a high-voltage electron beam that results in x-ray emission having energies characteristic of the elements present in the fiber. The amount of DCAA in the fibers was found to depend on the fiber residence time in the ethyl acetate coagulation bath. Removal of the DCAA is advantageous since residual DCAA not only slowly degrade the mechanical properties of polyaniline fiber but it is also a hazardous compound. Processes in our laboratory have shown that the dopant anions present in the as-spun polyaniline fibers can be replaced with more desirable dopant anions during (a) postspinning dopant manipulation, (b) the spinning process, and (c) fiber conversion to the insulating EB (dedoped) form as part of the spinning process, followed by redoping with another acid. The process for dedoping AMPSA=DCAA-doped fibers may be accomplished by immersion in deionized water, 0.1 M NH4OH, or steam extraction at 15 psi in an autoclave. Immersing the fibers in aqueous NH4OH is the fastest method and produces fibers having the lowest conductivity and the lowest dopant concentration. Higher hydroxide concentrations had only a marginal effect on the fiber composition. Steam extraction produces fibers with superior mechanical properties (see Table 2.8), but showed the slowest rate of AMPSA removal. TABLE 2.7 Variation in Properties of the As-Spun Polyaniline Fiber with Decreasing Residence Time in the Ethyl Acetate Coagulation Bath Coagulation Bath Residence Time (s)
Diameter (mm)
Density (g=cm3)
Denier (g=9000 m)
Conductivity (S=cm)
Tensile Strength (MPa)
Modulus (GPa)
Elongation (%)
115 36 17
44 + 2 45 + 2 43 + 2
1.49 1.47 1.57
20 21 21
790 + 110 790 + 110 960 + 110
142 + 6 130 + 3 149 + 4
4.1 + 0.3 3.2 + 0.8 4.5 + 0.6
8.0 + 1.3 8.9 + 3.6 5.7 + 0.5
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TABLE 2.8 Variation in Properties of the As-Spun Polyaniline Fiber after Steam Dedoping and Redoping with Methanesulfonic Acid Posttreatment As-spun Steam dedoped Redoped with CH3SO3H
Diameter (mm)
Density (g=cm3)
Denier (g=9000 m)
Conductivity (S=cm)
Tensile Strength (MPa)
Modulus (GPa)
Elongation (%)
43 + 2 31 + 2 36 + 2
1.57 1.24 1.65
21 8.4 15
960 + 110 160 + 40 1070 + 180
142 + 6 663 + 53 236 + 15
4.1 + 0.3 21 + 5 7.7 + 1.5
8.0 + 1.3 5.4 + 3.5 5.4 + 1.2
The AMPSA-doped polyaniline fiber that was exposed 17 s in ethyl acetate coagulation bath (Table 2.7) was subsequently exposed to steam at 20 psi for 2 h to dedope the fiber and was then reprotonated by soaking the dedoped fiber in 10 wt% methanesulfonic acid in methanol for 19 h. The methanesulfonic acid was dissolved in methanol instead of water, as this improves the tensile properties of the redoped fibers. The physical properties of these fibers (Table 2.8) are important in that they show that the steam-dedoped sample is the strongest PANI fiber made in our laboratory to date, in terms of tensile strength and modulus. Besides methanesulfonic acid, the fibers were redoped with a wide variety of acids by 16 h immersion in 1.0 M aqueous solutions of the acids listed in Table 2.9, which also lists the electrical and mechanical properties of those fibers. Dedoping the fiber causes the modulus and tensile strength to be higher than the as-spun polyaniline fiber, but results in a more brittle fiber. This result is similar to what was observed in Table 2.8 for the steam-dedoped fiber. Upon redoping the fiber, the lower the pKa of the redoping acid, the more conductive the fiber becomes. However, there appears to be no correlation between the pKa of the acid and the mechanical properties of the redoped fiber. All of the redoped fibers possessed a higher percent elongation than the EB fiber but had lower modulus and tensile strength. The fiber redoped with HCl showed the highest modulus whereas the fiber redoped with MSA possessed the highest percent elongation of the measured redoped fibers. Acrylic acid produced the fiber with the highest tensile strength of the measured redoped fibers. In order to effectively remove both AMPSA and DCAA and to eliminate the posttreatment dopantexchange stage, the possibility of adding a dopant-exchange process during the spinning process was explored. The feasibility of this approach was investigated using phosphoric acid. Phosphoric acid was chosen for this study because it imparts good thermal stability to the doped fiber and is a relatively inexpensive acid. The polyaniline fibers were produced from a 12 wt% PANI–AMPSA=DCAA dope solution that was extruded through a 100 mm diameter spinneret into an ethyl acetate coagulation bath. The residence time in the ethyl acetate coagulation bath was 38 s. The first godet was immersed in a 1.0 M aqueous phosphoric acid dopant-exchange solution. The residence time of the fiber in the first godet bath was 56 s. With the use of an empty first godet bath, control fibers were spun under the same conditions without dopant exchange. The stretch ratio between the two sets of godets was chosen to be TABLE 2.9 Variation in Properties of the As-Spun Polyaniline Fiber after Dedoping with Ammonium Hydroxide and Redoping with Various Acids PANI Fiber
pKa
As-spun EB fiber Redoping acid HCl CH3SO3H Oxalic acid Pyruvic acid Acrylic acid
2.2 2.0 1.23 2.39 4.25
Modulus (GPa)
Tensile Strength (MPa)
Elongation (%)
4.1 + 0.3 7.6 + 0.2
154 + 5 187 + 9
10.9 + 1.5 3.0 + 0.4
5.6 2.3 2.8 3.3 4.7
+ + + + +
0.7 0.2 0.5 0.1 0.9
156 104 106 107 179
+ + + + +
23 6 10 11 11
3.3 32.5 8.2 15.0 23.6
+ + + + +
1.0 5.6 2.9 1.4 4.8
Conductivity (S=cm) 424 + 22 <4 104 383 333 188 106 45
+ + + + +
17 13 12 21 9
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Conducting Polymer Fiber Production and Applications TABLE 2.10 Variation in Properties and Composition of the Polyaniline Fiber Spun with and without a H3PO4 Dopant Exchange Dopant Exchange
Diameter (mm)
Conductivity (S=cm)
Tensile Strength (MPa)
Modulus (GPa)
Elongation (%)
No Yes
72 + 2 70 + 4
300 + 10 315 + 40
55 + 2 69 + 2
1.6 + 0.3 2.1 + 0.2
20 + 15 26 + 5
Composition PANI–AMPSA0.58=DCAA0.18 PANI–H3PO4 0.70 AMPSA0.02=DCAA0.02
1.2:1, since it is expected that plasticization effect due to the presence of AMPSA in the polyaniline fiber that results in high stretch ratio would no longer exist. The stretching ratio limit between the two godets for these H3PO4 dopant-exchanged fibers was found to be 1.5:1. Table 2.10 illustrates the effect of including this dopant-exchange procedure on the composition, electrical, and mechanical properties of doped polyaniline fibers produced under the same spinning conditions. The composition of both as-spun polyaniline fibers were assessed using EDS. The 56 s residence time in the first godet bath containing the phosphoric acid solution resulted in substantially all of the AMPSA and DCAA in the solid fiber being replaced with phosphoric acid. The H3PO4 dopantexchanged fiber was found to have similar electrical and mechanical properties to the control AMPSAdoped polyaniline fiber. Residence times in the EA coagulation bath exceeding 38 s were subsequently found to produce weaker fibers. Residence times in the EA coagulation bath between 6 and 25 s gave a good combination of electrical and mechanical properties when a H3PO4 dopant-exchange step is employed. A residence time of 1 min in the dopant-exchange solution (at 1.0 M H3PO4) was found to be sufficient to replace AMPSA with phosphoric acid. Other acids are expected to be suitable for replacing the AMPSA and DCAA in a similar manner. The residence time in the coagulation bath was chosen such that acids having low pKa values will replace the AMPSA dopant molecules in the fiber. By contrast, the majority of DCAA molecules present in the fiber from the spinning solution were removed in the ethyl acetate coagulation bath. AMPSA has minimal solubility in ethyl acetate and is thus retained in the fiber. Since it is not possible to use a stretch ratio in excess of 1.5:1 between the two godets for increasing the mechanical properties of the H3PO4 dopant-exchanged polyaniline fibers, the effect of an additional thermal process to increase the fiber’s mechanical properties was explored. Phosphoric acid-doped polyaniline fiber spun from a 12 wt% PANI–AMPSA=DCAA dope solution was extruded through a 150 mm diameter spinneret into an ethyl acetate coagulation bath. The residence time in the ethyl acetate coagulation bath was 19 s. The first godet was immersed in a 1.0 M aqueous phosphoric acid dopantexchange solution with a residence time of 56 s. The stretch ratio between the two sets of godets was 1.25:1. In order to improve the mechanical properties, the fiber was left to dry under ambient conditions for 5 d. The physical properties of the as-spun fiber and the dried fiber are shown in Table 2.11. After remeasuring the mechanical properties, the fiber was heated in air at 708C for either 30 or 60 min and the properties were remeasured. As shown in Table 2.11, the mechanical properties of the thermally processed fiber are greatly enhanced, as the tensile strength increased from 84 to 195 MPa and the
TABLE 2.11 Variation in Properties of the Polyaniline Fiber Spun with and without a H3PO4 Dopant Exchange after Different Posttreatment Conditions Posttreatment
Diameter (mm)
As-spun After 5 d 30 min at 708C 60 min at 708C
70 70 70 70
+ + + +
2 2 2 2
Conductivity (S=cm) 460 440 415 375
+ + + +
25 25 25 20
Tensile Strength (MPa) 83 84 160 195
+ + + +
5 4 7 10
Modulus (GPa) 2.3 2.7 5.5 7.7
+ + + +
0.1 0.2 0.6 0.5
Elongation (%) 55 35 15 14
+ + + +
20 18 7 2
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TABLE 2.12 Variation in Properties and Composition of the Polyaniline Fiber Spun with a pH 10.8 Solution in the First Godet Bath Residence Time in the First Godet Bath (s)
Diameter (mm)
Tensile Strength (MPa)
Modulus (GPa)
Elongation (%)
Composition
60 15 3
52 + 2 60 + 2 84 + 2
204 + 7 167 + 6 43 + 3
7.6 + 0.3 5.7 + 1.2 1.3 + 0.2
3.2 + 0.2 8.2 + 5.4 51 + 27
PANI–AMPSA0.02=DCAA0.01 PANI–AMPSA0.08=DCAA0.05 PANI–AMPSA0.20=DCAA0.10
modulus increased from 2.3 to 7.7 GPa. Not surprisingly, this increase in strength came at the expense of the flexibility of the fiber since elongation at break was reduced from 35% to 14%. For acids such as triflic acid, the dopant-exchange process may not be economically or technically feasible. In an effort to completely remove DCAA and AMPSA from the fiber before it is collected on the bobbin and thus to minimize the amount of work at the posttreatment stage, we explored the feasibility of dedoping the polyaniline fiber into its EB oxidation state as part of the spinning process. This could be achieved by immersing the first godet in an ammonium hydroxide solution having a pH of 10.8. The polyaniline fibers were produced from a 12 wt% PANI–AMPSA=DCAA dope solution that was extruded through a 100 mm diameter spinneret into an ethyl acetate coagulation bath. The residence time in the ethyl acetate coagulation bath was 56 s. In order to verify the effect of base washing time on the mechanical properties, the residence time of the fiber in the ammonium hydroxide was varied. The heat tube was set to 808C and the second godet was rotated 1.2 times faster than the first godet followed by collection on a bobbin. The effect of residence time in the first godet bath on the fiber properties is summarized in Table 2.12. Based on EDS analysis of the as-spun fibers, a 3 s residence time in the first godet was found to be sufficient to remove approximately 90% of the DCAA and 65% of the AMPSA from the fiber. The fiber entering the first godet bath typically has a composition of PANI–AMPSA0.58=DCAA0.91. However, due to DCAA and coagulant EA trapped in the as-spun fiber, the fiber was soft and weak. As more solvent and coagulant were removed, the fiber was found to have a higher modulus but became more brittle. The conductivity of the fibers that resided in the first godet bath between 15 and 60 s, indicating that they were substantially dedoped, as their conductivity was less 104 S=cm. The fiber spun with a residence time of 3 s in the first godet bath has a conductivity of 10 + 4 S=cm. 2.4.2.4 Electronic Transport Properties of the Polyaniline Fibers Measuring the conductivity of the as-spun polyaniline fibers provides limited information regarding the effect of the processing conditions on the underlying structural order of these materials. The structural order of these materials can best be determined by measuring the electronic transport temperature dependence. For example, Monkman and coworkers [61] reported the temperature-dependent electrical conductivity of AMPSA-doped polyaniline fibers formed in an acetone coagulation bath after the fiber has been stretched to five times its original length. The conductivity of both the unstretched and stretched fibers showed a positive, then negative (insulator then metallic) temperature dependence of the polyaniline fiber as the temperature was increased from 15 to 300 K. It was found that the unstretched fiber had a higher ‘‘turnover’’ temperature than that of a thermally annealed AMPSAdoped polyaniline films. It should be noted that the value for the ‘‘turnover’’ temperature has been previously related to the crystallinity of the material [61], with lower ‘‘turnover temperatures’’ observed for more crystalline materials. This then implies that the AMPSA-doped polyaniline fibers are less crystalline than the corresponding films. The authors commented that it is presumably related to the method for producing the polyaniline fibers, in which the acetone coagulant rapidly removes the DCAA solvent from the fiber whereas during the production of evaporatively cast polyaniline films, the DCAA is removed slowly, which better allows for the formation of crystalline regions. Furthermore, the higher conductivity of the stretched polyaniline fiber indicates alignment of the polymer chains in the amorphous regions of the fiber upon drawing the fiber at 908C. The ‘‘turnover’’ temperature of
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Conductivity (S/cm)
the stretched polyaniline fiber was lower than 1000 2.7:1 that of the unstretched fiber, which is another indication of a more crystalline material. 800 In a previous study of polyaniline fibers 2.4:1 produced in-house, Bowman and Mattes [65] 2.1:1 reported on the electrical conductivity tem600 perature dependence of unstretched, AMPSA1.8:1 doped polyaniline fibers formed in an ethyl 1.5:1 400 acetate coagulation bath. Those fibers possessed a room temperature conductivity of 72 S=cm and their electrical conductivity 1.2:1 200 showed a positive (insulator) and then negative (metallic) temperature dependence as the temperature was increased from 4 to 330 K 0 0 50 100 150 200 250 300 350 with a turnover temperature of 260 K. While T (K) most other researchers employ postspinning stretching of the as-spun polyaniline fibers, FIGURE 2.25 Temperature dependence on the conductivin our improved spinning process stretching ity of the AMPSA-doped polyaniline fibers with different stretch ratios between the godet baths. was achieved during the spinning process by rotating the second godet pair faster than the first. The mechanical and conductivity properties of these fibers stretched at different ratios between 1.2:1 and 2.7:1 at 908C are described earlier in Table 2.6. The raw data for the temperature dependence of the conductivity of these polyaniline fibers is shown in Figure 2.25. Not surprisingly, Figure 2.25 shows that over the temperature range of 4 to 300 K, the conductivity of the polyaniline fibers increases when the stretching ratio between the godets is increased. The conductivity of these godet-stretched polyaniline fibers showed a positive (insulator) and then negative (metallic) temperature dependence as the temperature was increased from 4 to 330 K. The turnover temperature for the fibers spun with different stretch ratios between the godet bath is summarized in Table 2.13, and indicates that the turnover temperature (Tm) for all of these godet-stretched fibers (230–237 K) is greater than the turnover temperature measured by Bowman and Mattes [65] for the unstretched fiber (260 K). Structural disorder in ICP systems introduces localization phenomena and limits conductivity. The extent of disorder determines whether conductivity is metallic, insulating, or critical. In the classical definition, a metal should have a positive temperature coefficient of the volume resistivity (r), finite conductivity as T ! 0 K, and the logarithmic derivative of the temperature dependence of conductivity, i.e., the reduced activation energy (W), must show a positive coefficient. The reduced activation energy as a function of temperature is determined by W ¼
d ln s d ln r T dr ¼ ¼ d ln T d ln T r dT
(2:3)
TABLE 2.13 Parameters Calculated from the Experimental Data and the Heterogeneous Model Behavior for the AMPSA-Doped Polyaniline Fibers with Different Stretch Ratios between the Godet Baths Stretch Ratio 1.2:1 1.5:1 1.8:1 2.1:1 2.4:1 2.7:1
s (4 K) (S=cm)
s (300 K) (S=cm)
Tm (K)
rr
rm (mV cm)
rt (mV cm)
DE (meV)
s (A˚)
Lc (A˚)
31 70.4 90 80 122 165
252.5 479.5 562.0 622.0 761.0 871.0
237 234 230 235 230 230
8.15 6.81 6.24 7.78 6.24 5.28
10.23 5.00 4.84 4.12 3.50 3.10
3.08 1.65 1.39 1.25 1.03 0.90
4.10 3.67 3.62 4.07 3.65 3.52
5.8 6.3 6.9 6.1 6.8 7.5
59 60 66 56 50 76
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Reduced activation energy
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All these properties have been observed for different types of conducting polymers. Before deciding on modeling the conductivity, it is important to graph the reduced activation energy as a log–log plot of the temperature to determine in which regime (metallic, critical, or insulating) the conducting polymer lies. The temperature dependence of W in various regimes is as follows:
1
0.1
1.5
2
1.2:1
2.1:1
1.5:1
2.4:1
1.8:1
2.7:1
2.5
3
3.5 4 In (T )
4.5
5
5.5
FIGURE 2.26 Temperature dependence on the reduced activation energy of the AMPSA-doped polyaniline fibers with different stretch ratios between the godet baths.
1. In the insulating regime, W has a negative temperature coefficient (i.e., a positive slope). 2. In the critical regime, W is temperatureindependent for a wide range of temperatures (i.e., a zero slope). 3. In the metallic regime, W has a positive temperature coefficient (i.e., a negative slope).
Bowman and Mattes [65] also showed that the low-temperature (<30 K) transport properties indicate that the reduced activation energy had a positive temperature coefficient, which is indicative of this material lying on the metallic side of a disorder-induced metal–insulator phase boundary. The reduced activation energy for these stretched polyaniline fibers is shown in Figure 2.26. Since the slope is negative for all of the godet-stretched fibers, this indicates that these fibers are on the insulating side of the disorder-induced metal–insulator transition. This is in contrast to that observed for the unstretched polyaniline fiber, which lies on the metal side of the disorder-induced metal–insulator transition. A possible explanation for this behavior of producing a more conductive, but a more disordered polyaniline fiber is that stretching the fiber between the godets at 908C during the spinning process improves the metallic properties of the highly order domains whereas increasing the disorder in the amorphous regions that separate the ‘‘metallic islands.’’ An alternate explanation may be that mechanically stretching the fiber with a heat gradient between the godets increases the size of the highly order domains. This would make the localization lengths smaller while leading to an increased mean free path, and therefore cause an increase in the conductivity of the fiber in the direction of the stretch. Finally, it may be that stretching the fiber increases the size of both the amorphous and highly ordered domains, which permits these regions to migrate further apart. This would make the relative permeability finite and open the Coulomb gap in the density of states near the Fermi level. Thus, it may be that stretching opens a Coulomb gap and thus induces a transition from the metallic to insulating region. Since the stretched fibers were in the insulating regime, the localization length of these materials were determined by analyzing the experimental data between 4 and 100 K using the equation for Mott’s hopping [66]: 14 ! T0 s(T ) ¼ s0 exp T
(2:4)
In the above equation, T0 ¼ 18=(kBN(EF)Lc3), where Lc is the localization length. From the fitting parameter T0, the localization length for the different stretched AMPSA-doped polyaniline fibers is summarized in Table 2.13. It has been reported that the localization length for an unstretched AMPSAdoped polyaniline fiber is 179 A˚ [65]. In contrast, the localization length for the stretched fibers is between 56 and 70 A˚. This result indicates that the stretch alignment of the polymer chains at elevated temperatures does indeed decrease the localization length as mentioned above. To account for the complex morphology of these materials, the heterogeneous model [67–69] was applied to the data plotted in Figure 2.25. According to this model (Equation 2.5) for highly conducting
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samples, the conductivity is well described by a series combination of quasi-1D metallic resistivity (high conductivity along the chain direction) and between extended metallic regions: Tm T1 þ r1 exp s1 ¼ rm exp T T0 þ T
(2:5)
The first term with the coefficient, rm , is the highly anisotropic metallic term, with T m the temperature corresponding to the phonon wave vector that spans the Fermi surface. The second term is Sheng’s expression developed in the fluctuation-induced tunneling (FIT) model [70], where T0 and T1 are the Sheng’s parameters depending on the details of the tunneling. T0 is the temperature at which there is a significant contribution to the conductivity from the FIT process, T1 is the effective barrier height for regions between the metallic ‘‘islands.’’ The height of the potential barrier (in meV) can be calculated from DE ¼ kBT1, and based on the work published by Sixou et al. [71], the width of the barrier (h in A˚) can be calculated from T1 pffiffiffiffiffiffiffiffiffiffiffiffiffi 2s ¼ 2mDE : h T0 Very good fits of the experimental data shown in Figure 2.25 were obtained using the heterogenous model. The resulting parameters are summarized in Table 2.13. The change in the conductivity with stretching is an indication of two processes that might be occurring in the structure: alignment of the polymer chains and unfolding the individual chains, which facilitates the interchain and intrachain transport of the carriers, respectively. The decreasing values of both geometric ratios, rm and rt, is due to the geometric change of crystalline and amorphous regions in the fiber caused by increased stretching of the fiber between the godet baths. The improved interchain transfer is also confirmed by the decreased resistivity ratio (rr ¼ s300 K=s4 K) and potential barrier as the stretch ratio increases (Table 2.13). There is also consistent change in the height of the potential barriers as the stretch ratio increases, and although it is small, it plays an important role in increasing the conductivity of the stretched polyaniline fiber.
2.5
Other Methods for Fabricating of Inherently Conducting Polymer Fibers and Fabrics
2.5.1 Polyaniline Blend Fibers The concept of forming electrically conductive fibers from polymer blends of ICPs with insulating conventional polymers (both thermoplastic and thermoset) began soon after the first generation of ICP fibers was reported. The majority of these electrically conductive fibers have been produced using polyaniline due to its known solubility in a wide range of solvents that have been traditionally used in fiber spinning. Andreatta and coworkers [54,55] first showed that electrically conductive polyaniline fibers could be prepared from polymer blends dissolved in 96% sulfuric acid. They demonstrated that 5 wt% solutions of polyaniline and the rigid chain polymer poly(p-phenylene terephthalamide) (Kevlar) dissolved in 96% sulfuric acid could also be processed into electrically conductive fibers. The solution was extruded at room temperature and dry-jet wet-spun into chilled water. It was observed that the dried fibers were exceptionally smooth and had a shiny, metallic appearance. The electrical conductivity of these as-spun polyaniline blend fibers was determined to be 1 S=cm. The formation of electrically conductive fibers of polyaniline blended with poly(p-phenylene terephthalamide) processed from sulfuric acid was also reported by Hsu et al. [72]. A thoroughly mixed solution of 0.2 wt% emeraldine base and 17.6 wt% poly(p-phenylene terephthalamide) was dissolved in sulfuric acid and this dope solution was extruded at 808C through a 20-hole spinneret into a 18C water coagulation bath. The continuous 20 filament yarns that were collected on the bobbin were thoroughly washed with water, which probably led to partial deprotonation of the emeraldine salt. However, the
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as-spun fibers were green, indicating that the incorporated polyaniline was primarily in the emeraldine salt state. The diameter of the individual fibers that make up the yarn was 25 mm. It was also noted that if the fibers are immersed in ammonium hydroxide before drying, they turn blue immediately indicating rapid conversion of the emeraldine salt in the blend fiber to the emeraldine base form. The addition of a 1 wt% polyaniline loading to form the electrically conductive blend fiber resulted in an 18% reduction in the initial modulus of the poly(p-phenylene terephthalamide) fibers from 76 to 62 GPa. Although the reduction in the initial modulus is outside experimental error, the initial modulus of the composite fiber is still in the range of a high-performance polymer. In contrast, the reduction of the tensile strength at break was only 7% from 3.0 to 2.8 GPa. Besides blending polyaniline with poly(p-phenylene terephthalamide) to produce electrically conductive fibers, Zhang et al. [73] showed that conductive fibers could also be produced from polyaniline blended with nylon-11 (poly-v-aminoundecanoyle). The fibers were produced from a concentrated solution of emeraldine base and nylon-11 dissolved in sulfuric acid, and the dope solution was extruded through a 12-hole spinneret into 208C water coagulation bath. The fibers were then stretched in 658C water with a drawing ratio of 2.75:1. SEM and transmission electron micrographs (TEM) of the monofilament indicated pronounced phase separation in the blend. The morphology of polyaniline in the fibers was fibrillar, which is valuable for producing conducting channels. The electrical conductivity of the fibers ranged from 106 to 101 S=cm depending on the concentration of the polyaniline in the blend and the percolation threshold was 5 wt%. This percolation threshold is lower than that for carbon black-filled poly(ethylene terephthalate) fibers (18%). Similarly, Kinlen and Frushour [74,75] reported on the fabrication of a series of fibers that were spun from a blend of polyaniline doped with dimethylacetamide and polyacrylonitrile dissolved in dimethylacetamide. The amount of polyaniline in the polyacrylonitrile matrix ranged from 1 to 40 wt%. The dope solutions were extruded through a 50-hole spinneret into a 558C coagulation bath that contained a 50:50 mixture of water and dimethylacetamide. The fibers then passed through a series of cascading drawing baths, each of which contained water at 988C. The tensile strength, elongation, and modulus of these polyaniline blend fibers were found to be between 1.0 and 2.5 gpd, 28% and 35%, and 40 and 50 gpd, respectively. The conductivity of these polyaniline blend fibers ranged between 107 and 106 S=cm, which indicated that the fibers were partially dedoped as part of the spinning process. While these examples represent polyaniline blend fibers, produced using a solution-spinning process, Fryczkowski et al. [76] showed that melt-spun polyaniline blend fibers could also be fabricated. To produce the electrically conductive fibers, polyaniline doped with different amounts of dodecylbenzenesulfonic acid was blended with isotactic polypropylene. Since polyaniline doped with dodecylbenzenesulfonic acid is not thermoplastic, the fiber must be melt-spun at temperatures below 2008C to avoid cross-linking reactions and backbone degradation. This mixture has good processing properties and was able to be thermally processed into fibers at 1908C. The molten mixture was extruded through the spinneret into air with the fibers subsequently collected on a rotating bobbin. The fibers examined had different diameters because of the compositions of the polymer mixtures from which they had been molded, but generally ranged between 10 and 200 mm. When viewed with the naked eye, the melt-spun fibers extruded with no dodecylbenzenesulfonic acid appeared black. In contrast, the melt-spun fibers containing dodecylbenzenesulfonic acid appeared green, with the intensity of the color depending on the content of the acid. When processed with a low acid concentration, grains of polyaniline diffused in the fiber are visible throughout the entire fiber volume, i.e., both in the core and surface layer, which leads to surface defects. Upon increasing the acid concentration, the polyaniline grains became smaller and diffused more evenly in the fiber. Furthermore, the dodecylbenzenesulfonic acid concentration in the fiber affected the conductivity of the fiber, with the highest conductivity that was reported (106 S=cm) during processing at high acid concentrations. Similarly, Kim et al. [77] demonstrated that melt spinning of electrically conductive fibers is possible by blending doped the forms polyaniline and polypyrrole with either isotactic polypropylene or lowdensity polyethylene. Binary blends containing up to 40 wt% of the ICP were melt extruded into fibers at 1508C for low-density polyethylene, and at 2008C for isotactic polypropylene as the matrix polymers.
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The conductivity of the melt-spun fibers was low (1010 to 107 S=cm) due to the nonhomogenous morphology of the as-spun fiber. As was observed by Fryczkowski et al. [76], the doped polyaniline or polypyrrole particles were not completely dispersed in either the isotactic polypropylene or low-density polyethylene matrix but instead formed aggregates. It was suggested that the melt spinning of electrically conductive fibers containing ICPs should be modified to include several extended mixing steps to improve the homogeneity of the ICP in the melt-extruded fiber.
2.5.2 Fabrics Coated with Polypyrrole and Polyaniline One of the most widely used approaches for fabricating electrically conductive textiles from intractable ICPs is to use an in situ approach to polymerize a submicron thick coating of an ICP onto an existing textile substrate. In this in situ polymerization technique, the fabric is immersed in a solution containing the ICP monomer, an oxidant, and the desired dopant anion. This process is industrially applicable because it can be performed using standard textile dyeing equipment using aqueous solutions for both aniline and pyrrole. In contrast to polyaniline, the intractable nature of polypyrrole, which prevents it from its ability to be processed into textile fiber, has resulted in the fabrication and properties of polypyrrole-coated textiles have been more widely studied than for polyaniline-coated textiles. The in situ polymerization of pyrrole to form an electrically conductive textile was first reported by Kuhn and coworkers [1,78–80]. The process for forming the polypyrrole-coated textiles was based on the immersion of a fabric into an aqueous solution containing pyrrole, ferric chloride, or ammonium persulfate to initiate the polymerization reaction and usually a sulfonated dopant anion. Using dilute solutions of pyrrole (0.015–0.03 M), the polymerization reaction occurs on the surface of the fiber and leads to the formation of a precipitate in the bulk liquid phase. The selection of the dopant anion affects both the surface resistivity and thermal stability of the polypyrrole-coated textiles [1,78,79]. For example, significant thermal dedoping of chloride-doped polypyrrole was observed at moderately low temperatures (<808C). The addition of sulfonated dopant anions such as anthraquinone-2-sulfonate and naphthalene disulfonate was shown to produce polypyrrole-coated textiles with an order of magnitude than the surface resistivity, and with improved thermal stability. The surface resistivity of these polypyrrole-coated textile could be varied from 5V=& to 10 kV=& by controlling the polymerization time, the concentration of the reactants in the polymerization bath, and the type of anion used to dope the polypyrrole coating. Based on this research, the industrial fabrics division of Milliken and Company introduced to the market in the early 1990s, a range of polypyrrole-coated textile products under the trade name of Contex. This in situ approach has been used by other researchers to deposit polypyrrole onto a wide range of textile substrates including polyester [78,80–84], nylon [78,85,86], poly(ethylene terephthalate) [85,86], and glass fibers [1]. Besides using an in situ polymerization of pyrrole to form an electrically conductive textile, other researchers have adopted a two-step process. The major advantage of the two-step process is that it can be easily adapted into a continuous process for industrial applications. However, the structure of the polypyrrole may be different to that obtained from the in situ polymerization process. Several variations to the a two-step process include first immersing the textile support in a solution containing the oxidant and desired dopant anion and then exposing the impregnated textile to either pyrrole vapor [87–89] or pyrrole dissolved in an aliphatic solvent [90] to initiate the polymerization reaction. Alternately, the textile support may be first exposed to pyrrole vapor and then immersed into an aqueous solution containing the oxidant and desired dopant anion [91]. The in situ polymerization of aniline to form an electrically conductive textile was also first reported by Kuhn and coworkers [78,79] in which the polyaniline was deposited onto a textile substrate from an aqueous solution containing aniline, ammonium persulfate, hydrochloric acid, and either the disodium salt of 2,6-naphthalenedisulfonic acid or 1,3-benzenedisulfonic acid. In these studies, the polymerization conditions were controlled in order to deposit the polyaniline layer only onto the textile support with no polymer precipitating in the bulk liquid phase. This was accomplished by using dilute solutions of aniline (0.03 M). A subsequent study by Tzou and Gregory [92] on the deposition of polyaniline to nylon-6 fibers was focused on investigating the reaction kinetics of the chemically oxidative polymerization of aniline
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with and without the presence of a nylon fabric. It was found that the reaction proceeded faster in the presence of the nylon fabric. The in situ polymerization of aniline to form electrically conductive textiles has been reported using polyester [78,93], nylon [78,92,94,95], poly(ethylene terephthalate) [85,86], and Nomex [96] as the textile support. In a variation of this in situ polymerization process, Forveille and Olmedo [97] reported on the development of a two-step continuous process for depositing polyaniline onto textile substrates. This process consists of first immersing the textile substrate in an aqueous solution of aniline hydrochloride. After draining the excess monomer solution and drying the textile substrate in air at 608C, polymerization was initiated on the fabric surface by immersing it in an oxidant solution consisting of potassium dichromate in 2 M HCl. The excess oxidant solution was drained from the polyaniline-coated fabric and then washed with 2 M HCl before drying in air at 608C. The conductivity of the polyaniline layer produced using this process was 20 S=cm with a mean growth rate that reached 150 nm of polyaniline per processing cycle. Multiple processing cycles can be used to increase the thickness of the polyaniline coating. Using a completely different method for coating textile substrates with polyaniline, Kim et al. [77] showed that polyester yarns could be continuously coated with polyaniline by passing the yarn through a polyaniline-coating solution. The coating solution consisted of between 3 and 10 wt% polyaniline doped with dodecylbenzenesulfonic acid dissolved in xylene. The polyester yarns were withdrawn from the coating bath in the presence of a dry airflow along the coating surface to enhance solvent evaporation. It was observed that the electrical resistance per unit length of fiber of polyaniline-coated yarns decreased from 170 kV=cm to 66 V=cm as the concentration of doped polyaniline on the coating solution increased from 3 to 10 wt%.
2.6 2.6.1
Applications of Electrically Conductive Inherently Conducting Polymer Fibers and Fabrics Electrochemical Actuators
ICPs represent a unique class of materials that can be fabricated into electrochemical actuators that possess the unique combination of high-stress generation, lightweight, and low operational voltages [98–102]. The mechanism for actuation is primarily due to the reversible transport of ions [103,104] and solvent molecules [105] between the polymer and the electrolyte during electrochemical oxidation and reduction. The resulting large dimensional changes generate stresses that exceed those of mammalian muscle (0.1–0.5 MPa) by at least an order of magnitude with less than 1 V driving voltage [100]. Additionally, linear strains up to 10-fold greater than piezoelectric polymers (typically <0.1%) are realized. To fabricate all-polymer ‘‘dry’’ actuator devices, the surrounding liquid electrolyte needs to be eliminated, which is typically accomplished by using a gel electrolyte. These solid-state actuators have the advantage of free-standing and can be operated in air. De Rossi and coworkers [106,107] first reported the development of solid-state linear actuators that makes use of polyaniline fibers as the actuating electrode. The polyaniline fibers were produced using a similar method to the one reported by Tzou and Gregory [29] in which the polyaniline fiber is produced by extruding a solution containing 20–25 wt% emeraldine base dissolved in DMPU through a spinneret into a water coagulation bath. The solid-state actuator was fabricated from polyaniline fiber that had been doped with perchloric acid and embedded in a solid polymer electrolyte. A metallic counterelectrode was fabricated by winding a 100 mm diameter copper wire along a helix to the external surface of the solid polymer electrolyte jacket. The actuator was operated by applying a voltage between the polyaniline fiber and the copper wire electrode. In this device configuration, no reference electrode was employed. Although the strain generated by the polyaniline fiber was relatively modest (0.2%–0.3%), the stress values developed by the dry actuator (2–3 MPa) are very interesting since they are approximately 10 times higher than that of human muscular system. In a slight modification to this solid-state actuator configuration, Lu et al. [108] reported the fabrication and performance of solid-state electrochemical linear actuators with a unique polyaniline solid-in-hollow fiber configuration. The basic device structure consists of a polyaniline solid fiber
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actuation electrode threaded into a polyaniline hollow fiber counterelectrode. A polymer gel electrolyte was used to separate the two electrodes and to serve as the ion source for the polymer electrodes upon redox cycling. The as-spun polyaniline AMPSA-doped solid and hollow fibers were redoped with triflic acid and this yielded greater electroactivity and stronger actuation in the gel electrolyte than the as-spun AMPSA polyaniline fibers. The actuator was operated by applying a voltage between the solid and the hollow fiber electrodes. In this device configuration, no reference electrode was employed. Isotonic strains of 0.9% and isometric stresses of 0.9 MPa were realized, which also exceeds that of skeletal muscle. While most of the research on ICP actuators has been carried out in aqueous or organic electrolytes, these systems suffer from narrow electrochemical potential windows and high volatility. These factors limit the device lifetime and performance. Lu and coworkers [109,110] investigated the electrochemical actuation of polyaniline fibers in the room temperature ionic liquid 1-butyl-3-methyl imidazolium tetrafluoroborate ([BMIM]þ [BF4]). Room temperature ionic liquids are ideal for the fabrication of ICP actuators since they possess high ionic conductivity, large electrochemical windows, excellent thermal and electrochemical stabilities, and negligible evaporation. When using [BMIM]þ [BF4] as the electrolyte, it was found that the electrochemical actuation of polyaniline fibers was strongly affected by the dopant anion. The as-spun polyaniline AMPSA-doped fibers showed negligible electroactivity and actuation in the [BMIM]þ [BF4] electrolyte. However, by substituting triflic acid as the dopant, good electroactivity and actuation in the [BMIM]þ [BF4] electrolyte were obtained. The linear actuation performance of the polyaniline fiber was measured by isotonic strain during one redox cycle reached a maximum value of 0.28% (applied load of 1 g) whereas the isometric stress was 1.8 MPa. In lifetime tests of the fiber, both electroactivity and electromechanical actuation continued without significant decrease (<10%) in either stress or strain after 1 million cycles. Lu et al. [111] later reported on the development of solid-state electrochemical linear actuators with a polyaniline yarn-in-hollow fiber configuration using an ionic liquid electrolyte. These yarn-in-hollow fiber actuators, which were constructed using a triflic acid-doped polyaniline solid fiber inserted into a triflic acid-doped polyaniline hollow fiber. A porous polyacrylonitrile insert separated the two electrodes, which contained the [BMIM]þ [BF4] electrolyte (Figure 2.27). It was demonstrated
Yarn
Separator
Hollow fiber
Acc.V Spot Magn 10.0 kV 3.0 35x
Det WD SE 26.5 Hivac
1 mm
FIGURE 2.27 SEM image of yarn-in-fiber electrochemical linear actuator fabricated by threading eight polyaniline yarns into a polyaniline hollow counterelectrode with a porous polyacrylonitrile separator inserted between the two electrodes containing the ionic liquid electrolyte.
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that a 20 monofilament polyaniline yarn with twist ratio of 1 turn per inch showed even larger strain (0.82%) and stronger force generation (0.102 N) than a bundle of 20 untwisted fibers (0.80% and 0.066 N). It is well known that there is an optimum twist ratio for yarns that leads to higher tensile strength compared to the same number of untwisted fibers. This translates into higher force generation when the yarn is used as the actuating electrode. Stress generation of these actuators was 0.42–0.85 MPa. Depending on the weight of the object, this yarn-in-fiber configuration would allow the combination of an appropriate number of yarns as the actuation electrode to accomplish the mechanical task.
2.6.2
Vapor and Humidity Sensors
Real-time vapor detection and monitoring have become increasingly important to prevent human exposure to toxic vapors that may be encountered in both industrial and battlefield environment applications. Vapor detection using ICPs have been evaluated using microelectronic devices such as chemi-resistors (interdigited array transducers) [112,113] and field-effect transistors (FETs) [114,115]. The chemi-resistive behavior arises from the absorption of the vapor, which interacts with the polymer backbone or the dopant molecule and thereby modulates the carrier mobility and number of free carrier charges available. In order to detect, quantify, and discriminate among various vapors, arrays of chemiresistive sensors are fabricated into a microelectronic device that is commonly referred to as an ‘‘electronic nose.’’ In this approach, classification and identification of the vapor is not achieved primarily through use of a ‘‘lock-and-key’’ design in which a detector responds in a highly selective fashion to a specific target compound or class of target compounds. Instead, each detector responds to a broad class of stimuli, with the collective response of the many different ICP sensors in the array providing a unique fingerprint for each vapor of interest. In contrast to vapor detection using handheld detectors fabricated from interdigited arrays, which can only provide point source detection, chemical sensing using a smart fabric offers the opportunity of using the large surface area of the fabric for improved selectivity and expanded dynamic response and spatially distributed sensing. The concept of an active sensing ICP fabric could eventually be translated into a wearable, lightweight, integrated sensor protection garment. Collins and Buckley [85,86] first demonstrated that polyaniline and polypyrrole-coated poly(ethylene terephthalate) or nylon fabrics prepared using the in situ polymerization approach could behave as a chemi-resistive sensor for detecting a wide variety of vapors. It was reported that these polyaniline- and polypyrrole-coated fabrics have detection limits in parts per million when exposed to the industrial pollutants nitrogen dioxide and ammonia as well as to the chemical warfare-simulant dimethyl methylphosphonate. Not surprisingly, it was also shown that the resistance of the conductive fabric was affected by the dopant incorporated into the polypyrrole, but also the type of ICP deposited onto the textile support. For example, the polyester fabric coated with polypyrrole doped with naphthalene disulfonate possessed a greater sensitivity for the detection of ammonia and nitrogen dioxide than for the fabric coated with polypyrrole doped with anthraquinone-2-sulfonate. However, the polyaniline-coated fabric possessed greater sensitivity for the detection of ammonia and nitrogen dioxide than the polypyrrole-coated fabrics. It was observed that polyaniline- and polypyrrole-coated fabric chemiresistive sensors suffer from a drifting baseline that is associated with the adsorption of water from the environment as the humidity changes. Kincal et al. [81] later reported on the chemi-resistive behavior of polypyrrole-coated fabrics containing different dopant anions when exposed to ammonia and hydrochloric acid vapors. The polypyrrole-coated fabrics were prepared by depositing a polypyrrole layer onto polyester fabrics using the in situ polymerization approach. Initially, an order of magnitude reversible change in conductivity was observed when these fabrics were sequentially exposed to ammonia and hydrochloric acid vapors but the magnitude of the change decreased with time. In order to probe the sensitivity of the polypyrrole-coated textiles to less reactive gases, the HCl was replaced with carbon
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Resistance (Ω)
dioxide and a twofold change in conductivity was observed. In this case, although the magnitude of the conductivity change was smaller, it was more highly reversible and did not exhibit hysteresis. However, the authors did not report the detection limit of these fabrics to the different vapors included in this study. While these studies have shown that polyaniline- and polypyrrole-coated fabrics can behave as chemiresistive sensors, it is also expected that similar chemi-resistive behavior will be achieved using ICP monofilaments and yarns, which are woven or stitched into an existing fabric. Ultimately, vapor detection fabric will need to comprise of an array of sensors with the patterned response fed into a microprocessor in order to identify and determine the concentration of the vapor. The relatively low cost of production and subsequent processing of these electrically conductive fabrics provide the opportunity for designing a textilebased ‘‘electronic nose,’’ which possess built-in massive redundancy for the sensor array so that device will still function even if there is a localized break (rip, tear hole, etc.) in the fabric network. Specifically, in case some area of the sensor system fails, there is a massive redundancy in the garment such that other areas may still be able to acquire the information. This will greatly increase the reliability of the system. To demonstrate the concept of using polyaniline monofilaments as a chemi-resistive sensor, our laboratory has explored the humidity-dependent chemi-resistive behavior of an AMPSA-doped polyaniline monofilament. It has been well documented that the electrical conductivity of doped polyaniline films and powders changes with the humidity due to the absorption of water into the polymer [116]. The water that is absorbed into the polymer takes part in the charge transfer, leading to a decrease in the resistance of the polymer. This humidity-dependent resistivity change is not due to a chemical reaction between water and the polymer backbone, and hence the presence of water will not degrade the p-conjugated nature or polyaniline backbone. Since, the resistivity changes induced by water adsorption have a high degree of reversibility, humidity sensors based on polyaniline films have been successfully fabricated [117–120]. Figure 2.28 shows the resistance change of a polyaniline monofilament as a function of relative humidity at room temperature (26.58C + 0.58C). The resistance measurements in Figure 2.28 were taken when the change in resistance at a particular relative humidity is less than 1% (approximately at equilibrium). Although often overlooked in the fabrication and testing of polyaniline humidity sensors, the resistance of polyaniline changes with both temperature and humidity. It is therefore imperative that the resistance signal acquired from the humidity sensor fabric is decoupled by using two distinct types of polyaniline fibers. The temperature is obtained by encapsulating one of the polyaniline fibers with a water-impermeable material, such that any change in resistance of this fiber is only associated with changes with temperature. The other polyaniline fibers are directly exposed to ambient air and there4000 fore changes in the resistance of these fibers are both temperature- and humidity-dependent. The differ3500 ence in the resistance signals between the separated temperature and humidity signals is amplified by a 3000 differential amplifier. A fabric was subsequently devel2500 oped in our laboratory in which a combined temperature and humidity sensor fabricated from polyaniline 2000 fibers is embedded within the textile [121]. 1500
2.6.3 Strain Sensors It was found by De Rossi and coworkers [107,122] that conventional fabrics that have been coated with a thin ICP layer possess the remarkable property of piezoresistivity. In other words, resistance of these fabrics varies with the application of a force to the
1000 10 20 30 40 50 60 70 80 90 Relative humidity (% RH)
FIGURE 2.28 Resistance of a 136 mm long AMPSA-doped polyaniline fiber as a function of percent relative humidity at room temperature (26.58C + 0.58C).
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fabrics that gives rise to a dimensional change along its length. These studies used a woven Lycra=cotton fabric coated with perchlorate-doped polypyrrole using the in situ polymerization technique developed by Kuhn and coworkers [1,78–80] and possessed an unstressed surface resistivity of 3V=&. The piezoresistive property of the polypyrrole-coated Lycra=cotton fabric was quantified by exerting successively small increments of uniaxial stretching along the two orthogonal directions in the plane of the fabric. As the degreeof extension increased, the resistance of the fabric decreased. The strain gauge factor, GF ¼ (DR=R Dl=l) of this first hypothesis, of a 34 mm thick polypyrrole film was measured using the same methodology as for the polypyrrole-coated fabrics, and it was found that the GF ranged from 0.45 to 0.9 depending on the degree of doping. Comparing this value with that of the polypyrrole-coated Lycra=cotton fabric, the polypyrrole film was much smaller and opposite in sign. Therefore, the piezoresistivity of the fabric must be due to more than the polypyrrole. Since the piezoresistivity of carbon fiber felts under tension and compression are similar to that of the polypyrrole-coated Lycra=cotton fabric, the authors proposed that the piezoresistive properties of the polypyrrole-coated Lycra=cotton fabric originate from a progressive increase in the number of contacts, the area contacting the microfibers, and finally the bundles themselves as the applied load increases. The piezoresistive properties of the polypyrrole-coated Lycra=cotton fabric has been successfully exploited in the fabrication of comfortable, wearable devices able to record and control human posture and gesture. Monitoring body kinematics and analyzing posture and gesture is an area of major importance in bioengineering and other related disciplines. Such techniques can be used for injury prevention, rehabilitation, sports technique modification, and medical treatment. For example, De Rossi and coworkers [107,122] reported on the development of a sensing glove, which was fabricated by securing strips of the polypyrrole-coated Lycra=cotton fabric on the upper side of the each finger of a silk glove. Each strain sensor passed its resistance signal to a different channel of the data acquisition set in order to detect flexion of each finger. De Rossi and coworkers [122] also reported on the construction of a sensorized leotard to monitor trunk and upper limb position and motion by attaching four strain sensors to the garment. Similarly, Wallace and coworkers [83] reported on the piezoresistive properties of a Lycra=cotton fabric which was coated with a thin layer of 1,5-naphthalenedisulfonate-doped polypyrrole using the in situ polymerization technique. A strip of this polypyrrole-coated Lycra=nylon fabric was then utilized in the fabrication of an intelligent knee sleeve in which different audio signals are emitted based on the strain detected by the fabric sensor. While the piezoresistive effect has been observed for polypyrrole-coated fabrics, our laboratory 0 has explored the piezoresistive properties of an electrically conductive fabric in which AMPSA−2 doped Panion monofilaments were knitted (see −4 Figure 2.22c) into a woven nylon support fabric [121]. The piezoresistive property of this fabric −6 was quantified by exerting successively small increments of uniaxial stretching. The change in the −8 resistance of the nylon fabric upon stretching is −10 shown in Figure 2.29 and indicates that the −12 resistance of the fabric decreased as the strain increased. This phenomenon is similar to what −14 was observed for the polypyrrole-coated fabrics by other researchers [83,107,122]. The change in −16 40 50 60 70 0 10 20 30 resistance of this fabric appears to fall within two Strain (%) distinct regions. When the strain is less than 20%, increasing the strain has a minimal effect on FIGURE 2.29 Changes in the resistance of the polythe resistance of the fabric. This is presumably aniline monofilaments were knitted into a woven nylon support fabric when stretched from 0% to 60%. due to the polyaniline fibers knitted in the fabric
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gradually becoming taut. At strains greater than 20%, the data shown in Figure 2.29 suggests a linear relationship between the resistance of the fabric and the strain. In this region, the lowering resistance can be ascribed to increased alignment of the polymer chains in the polyaniline monofilaments.
2.6.4 Resistive Heaters
Temperature change ∆T (K)
Resistive heating fabrics are traditionally fabricated using either stainless steel or carbon fibers embedded in an existing fabric. When such a conductive textile is supplied with an appropriate voltage it heats by Joule effect. Regardless of the conductive material, resistive heating fabrics are well suited for the large area radiant or direct contact heating. For example, several studies have found that using carbon fiber resistive heating blankets was more effective in treating hypothermia than either infrared reflective covers (space blankets) or circulating water mattresses and comparable in performance to forced air warmers [123–125]. The surface resistivity of ICP fabrics also makes these materials suitable for general resistive heating applications, such as heated clothing and car seats, as well as electric heating in floors and walls. Furthermore, ICP fabrics allow power to flow directly through the fabric without incorporating embedded fabric wiring. The continuous conducting fabric also produces a more even heat flow and is more pliable than fabrics containing wires. Jolly and coworkers [126,127] first described the fabrication and properties of polypyrrole-coated textiles for building heating panels, heating sheets for car seats, heated winter sportswear, and heated gloves for medical applications. When compared to other resistive heating devices, it was found that the polypyrrole-coated textiles exhibited improved temperature homogeneity, low power requirements, and could easily be cut and sewed onto different substrates for a large range of applications. It was noted that the power density required for current building heating panels must be greater than 150 W=m2 whereas the polypyrrole-coated textiles showed higher power densities between 300 and 400 W=m2. Kaynak and coworkers [82,128] later reported on the effect of different dopant anions incorporated in the polypyrrole-coated textiles on the heat generation of these materials. The polypyrrole layer was deposited onto a polyester=Lycra fabric using the in situ polymerization approach and was doped with anthraquinone-2-sulfonate, naphthalene-2-sulfonate, p-toluenesulfonate, or perchlorate. At an applied voltage of 24 V, the polypyrrole-coated fabrics, from all the four different dopant systems showed an increase in temperature with the anthraquinone-2-sulfonate-doped polypyrrole coating the most effective heat generator (DT 208C) whereas the sodium perchlorate dopant system was the least effective (DT 38C). The power density per unit area achieved in the anthraquinone-2-sulfonatedoped polypyrrole-coated fabric was 430 W=m2, 200 W=m2 for naphthalene-2-sulfonate, 150 W=m2 for p-toluenesulfonate, and 55 W=m2 for perchlorate, respectively. 30 While all previous ICP fabrics used in resistive heating studies have used polypyrrole-coated fab25 rics, our laboratory has explored the heat gener20 ation of Panion monofilaments [129]. The thermal characteristics of an AMPSA-doped polyaniline 15 fiber were investigated under ambient conditions by applying a constant voltage of 4.5 V. For these 10 experiments, the polyaniline fiber was 8.5 mm long and had a diameter of 95 mm. Figure 2.30 5 shows the stable temperature change as a function of applied voltage for the conductive poly0 3 4 5 0 1 2 aniline fiber under ambient conditions, where the Voltage (V) baseline temperature was 238C. The time to reach a stable temperature varied from 10 s for a temFIGURE 2.30 Steady-state temperature change as a perature of only a few degrees up to 40 s for a function of applied voltage for an AMPSA-doped polyaniline fiber. 208C temperature change. Figure 2.30 shows that
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the temperature change is proportional to the square of the voltage on the fibers. This is not an unexpected result as, according to Ohm’s law, the 10000 electrical energy supplied to the fiber is given by 8000 Q ¼ P t ¼ V2=R t whereas according to heat transfer theory, the thermal energy is proportional 6000 to the temperature change (Q / DT). If we combine these two expressions since electrical energy 4000 supplied to the fiber is equal to the thermal energy released by the fiber, then DT / V2. It should be 2000 noted that the raise in ambient temperature produced by the polyaniline fiber is essentially identi0 cal to that for the polypyrrole composite fabrics 0 20 40 60 80 100 120 [82, 126–128] with a DT between 208C and 258C. Time (s) An interesting phenomenon observed in this FIGURE 2.31 Resistance of an AMPSA-doped polystudy was that applying an overloading current aniline fiber as function of time when an overload or voltage was found to irreversibly destroy the voltage was applied to the fiber. conductive nature of the polyaniline fiber. The resistivity of an AMPSA-doped polyaniline fiber as a function of time is shown in Figure 2.31 for the situation where an overload voltage of greater than 4.5 V was passed through a polyaniline fiber with a length of 12 mm and a 95 mm diameter. The temperature measured upon the resistance jump was 488C (DT of 258C). This is in contrast to the thermal decomposition temperature for AMPSA-doped polyaniline fibers, which is 1808C as determined by thermogravimetric analysis. Similarly, after this polyaniline fiber was redoped with triflic acid and HCl, the temperature at which this fuse-like behavior occurred was between 388C and 968C, respectively. This voltage or current overloading is a different phenomenon than found in most conductive wires, which can only be destroyed by melting. It was also observed that when the electrical conductivity of the fiber has been substantially destroyed, the structural integrity of the fiber is preserved. It is believed that conductive polymer fibers are destroyed by the alteration of their conjugated structures but that this occurs at temperatures below the temperature at which dopants within the fiber are lost (typically between 1008C and 2508C depending on the dopant anion) or at which the polymer backbone decomposes (>3008C). The low electrical destruction temperature makes the use of these materials suitable as fuses and safe resistive heating devices such as a heating element placed in the vicinity of the skin. Fiber resistance (Ω)
12000
2.6.5
Electromagnetic Interference Shielding
Electromagnetic interference (EMI) is one of the unfortunate by-products of the rapid proliferation of electronic products and telecommunication equipment. EMI shielding is not a new concern due to its interference with other electronic devices but the need has become more widespread due to continuing miniaturization, increasing sensitivity, and importance of these electronic devices. The traditional approach for EMI shielding has been the use of metals as they have high electrical conductivity and dielectric constants. The disadvantage of using metals either in polymer-filled composites or as coatings is that they have poor corrosion resistance and impose severe weight penalties, especially in aerospace systems. Carbon has also been used in EMI shielding applications as a conductive filler in composite materials due to its high electrical conductivity, good chemical resistance, and low density. ICPs alone, and when prepared as composite materials, are becoming increasingly attractive for EMI shielding applications due to their higher electrical conductivity than graphitized carbon-based composites, low density, ease of processing, and finally their unique absorption shielding mechanism. This EMI shielding mechanism differs from that of metals and carbon, which is based on reflection, and it is preferred by the military for stealth and camouflage applications [130]. Consequently, the EMI shielding characteristics of both polyaniline- and polypyrrole-coated fabrics have been widely studied. For these ICP-coated
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fabrics to be used in EMI shielding, they must possess a high EMI shielding efficiency (SE), defined as the attenuation of an electromagnetic wave produced by its passage through a shield. It is measured as the ratio of the shield strength before and after attenuation at a specific frequency. It should be noted that the shielding efficiency for EMI control is largely a function of surface conductivity of the material. For most industrial applications, a shielding efficiency of 30 dB is considered a useful attenuation value because it will prevent passage of 99.9% of the EMI. In one of the first reports on using an ICP-coated fabric for EMI shielding, Trivedi and Dhawan [131] measured the EMI SE characteristics of polyaniline-coated fabrics using the coaxial transmission line method in the frequency range of 1000 kHz to 1 GHz. The polyaniline-coated fabrics were prepared using an in situ polymerization technique and contained different dopant anions (5-sulfosalicylic acid, p-toluenesulfonic acid, benzenesulfonic acid, or 4-hydroxybenzenesulfonic acid). It was found that at higher frequencies (0.1MHz to 1 GHz), the SE is between 16 and 18 dB whereas at lower frequencies, it is greater than 40 dB. The authors observed that the doping level and type of dopant, as well as the thickness of the PANI layer, had strong effect on the SE of the PANI-coated fabrics. Dhawan et al. [132] later evaluated the performance of polyaniline-coated fabrics, including polyester and silica fabric for the control of EMI shielding at 101 GHz. EMI shielding measurements revealed that the SE of these coated fabrics increased with the thickness of the polyaniline layer. For instance, the HCldoped polyaniline-coated polyester fabric, with a resistivity of 160 V cm, showed an SE of 6.3 dB. However, with doubled and tripled thickness, the SE of coated fabric respectively increased to 25.52 and 29.1 dB. Similarly for the p-toluenesulfonic acid-doped polyaniline-coated polyester fabric (resistivity of 40 V cm), the SEs increased from 17.77 to 47.63 dB when the thickness of the polyaniline layer was increased threefold. Moreover, the polyaniline-coated silica fabric possessed better EMI shielding than the polyester-coated fabric when prepared under identical conditions. An SE value of 35.61 dB was obtained for the coated silica fabric whereas the SE for the coated polyester fabric was 21.48 dB. It was speculated that this increase in performance was due to the fact that the silicon fabric retained more polyaniline in the interstices of the fabric than the polyester fabric. As a continuation of this study, Dhawan et al. [132] investigated the EMI response of these polyaniline-coated fabrics in the microwave range, W-band, RFI range, and UV–vis NIR range. These polyaniline-coated fabrics exhibited an SE of the order of 30–40 dB over the range of 100–1000 MHz as measured by coaxial transmission line method and an SE of 3 to 11 dB in the range 8–12 GHz as showed by the microwave absorption studies. The reflectance and transmittance studies showed that in the UV–vis NIR range, 98% of the energy was absorbed by the polyaniline-coated fabrics, with polypyrrole-coated fabrics (95%), and only 2% was reflected. Kim et al. [133] investigated the EMI shielding performance in 50 MHz to 13.5 GHz frequency range of polypyrrole-coated nylon fabrics. In this study, polypyrrole was electrochemically deposited onto the polypyrrole-coated nylon fabric prepared using the in situ chemical polymerization approach to increase the thermal stability and conductivity of the polypyrrole-coated fabric. It was found that the multilayer polypyrrole-coated fabric was better than chemically deposited polypyrrole-coated fabric for EMI shielding. The values for SE were in the range of 5–40 dB depending on the conductivity of the polypyrrole-coated fabric.
2.7
Electrospun Fibers of Inherently Conducting Polymers
Semiconducting one-dimensional (1D) nanofibers or nanowires are of interest for a wide variety of applications including interconnects, functional devices, and molecular sensors as well as for fundamental physics studies. Devices have been fabricated from semiconductor, and carbon nanotubes, and more recently from ICP nanofibers. It has been predicted that ICP nanofibers will have unique electrical, optical, and magnetic properties [134]. Several different methods for producing these ICP nanofibers have been developed with or without the aid of a template. The template-based methods involve synthesizing a tubular structure of the ICP within the pores of a support membrane, such as an alumina membrane [135] or a track-etched polycarbonate membrane [136]. However, more recent work has
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focused on the development of template-free synthetic routes for fabrication of ICP nanofibers. For example, an in situ doping polymerization approach has been described in which the use of large organic anions results in polyaniline nanofibers and nanotubes with diameters around 650 nm [137]. An interfacial polymerization method was reported by Huang and Kaner [138], in which polyaniline nanofibers were produced at the interface of two immiscible liquids. These polyaniline nanofibers had nearly uniform diameters between 30 and 50 nm. Similarly, Manohar and coworkers [139] developed an approach where polyaniline nanofibers could be fabricated by seeding a conventional chemical oxidative polymerization of aniline with very small amounts of biological, inorganic, or organic nanofibers, which changed the morphology of the resulting doped polyaniline powder from nonfibrillar to almost exclusively nanofibers. For certain applications, it is highly desirable to have the lengths of the ICP nanofibers to be greater several hundred micrometers, which is significantly longer than lengths of the ICP nanofibers prepared using the aforementioned methods that typically vary from 500 nm to several micrometers. Under appropriate conditions, ICP nanofibers with lengths that are several meters long can be produced using an electrostatic, nonmechanical process called electrospinning. Formhals [140] first patented the concept of producing polymeric nanofibers using an electrospinning process in the 1930s. The electrospinning process has received much attention in the last decade since diameters of electrospun fibers are at least one to three orders of magnitude smaller than fibers made by conventional spinning techniques, which use mechanical forces to produce fibers by extruding polymer fluids followed by subsequent drawing of the resulting monofilaments. The main limitation of preparing ICP nanofibers by the electrospinning technique has been the difficulty of preparing stable, highly concentrated polymer solutions. This is the same problem that has generally prevented the fabrication of macrosized ICP fibers by conventional wet-spinning techniques. Although the method for extruding the fibers is significantly different (electrostatic vs. mechanical forces), the required physical properties of the polymer solution are similar. However, other techniques developed for fabricating ICP fibers and textiles have been successfully adopted to fabricate ICP nanofibers. These techniques include blending the ICP with known fiber-forming polymers to obtain high-viscosity spinning solutions or coating the exterior of electrospun nanofibers with an ICP layer via an in situ polymerization process.
2.7.1
Electrospinning Process
The electrospinning process is a variation of the better known and understood electrospraying technique. In the electrospinning process, a high electric field is generated between a highly viscous polymer solution held by its surface tension at the end of a capillary tube and a metallic target, as shown in Figure 2.32. As the intensity of the electrical field increases, the surface of the liquid hemispherical drop, suspended at equilibrium at the capillary tip, elongates to form a conical shape, known as the Taylor cone [141]. The ‘‘balancing’’ of the repulsive electrostatic force with the surface tension of the liquid causes this distortion. When the voltage reaches a critical value (0.5 kV=cm), the charge overcomes the surface tension of the deformed drop and a jet is produced. As a result of the low surface tension in lowviscosity solutions, the jet breaks apart into a series of droplets. This is the basis of electrospray technology. As the jet travels from the anode to the cathode, rapid evaporation of solvent molecules occurs that reduces the diameter of the jet in a ‘‘cone-shaped’’ volume. This is called the ‘‘envelope’’ cone (Figure 2.32). The as-spun dry fibers accumulate on the surface of the collection screen. This process results in a porous, nonwoven mesh of nanofibers. Polymer melts have also been processed into nanofibers, and the metallic target used to quench the as-formed molten fiber mats. Replacing the metal target with a grounded coagulation bath target leads to instantaneous demixing of the polymer solution, which leads to the production of continuous nanofiber filaments from solvents with high boiling points. The electrospinning process has explored the types of polymer and solvent systems from which fibers can be produced. In these studies, more than 30 different types of both conventional and specialty polymers that include polyethylene oxide [142–147], nylon [148,149], polyaramid [150]
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Metal target
Fluid droplet
Envelop cone
High voltage
FIGURE 2.32
Schematic of the electrospinning process.
and polyacrylonitrile [151], DNA [152], poly(L-lactic acid) [153], silklike proteins [154,155], and elastin mimetic peptides [156] have successfully been electrospun into nanofibers. Early work in the field of electrospinning was limited to fibers with average diameters in the 500– 2000 nm range. Currently, the diameter electrospun fibers fall in the range of 50–500 nm, which is due to advances in understanding the electrospinning process itself. Moreover, Fong and Reneker [157] reported the preparation of nanofibers of a styrene–butadiene–styrene block copolymer that have diameters as small as 3 nm (i.e., approximately six polymer molecules across the diameter of the fiber). Many of these studies have simultaneously addressed some of the processing or fiber diameter and morphology relationship. The processing parameters that have been considered include solution concentration and viscosity [142,145,147,157–161], surface tension [142,145,147], and electric field strength [144,157,160]. The most studied parameter has been the effect of polymer concentration and resultant solution viscosity, which has been found to affect the fiber diameter [142,145,147,160], initiating droplet shape at the capillary tip [158] and the jet trajectory [142,159,161]. For the electrospinning process to yield nanofibers, it is essential to obtain a viscosity of the polymer solution between 100 and 4000 cP, although the majority solutions possess a viscosity in the range of 1000–2000 cP. At low viscosities (<100 cP), surface tension is the dominant influence and below a certain concentration, the jet breaks apart to form drops instead of fibers. At high concentrations (>4000 cP), nanofiber production is prohibited by the inability to control and maintain the flow of the polymer solution through the capillary and by the cohesive nature of the highly viscous solution. Furthermore, increasing solution viscosity has been associated with production of fibers with larger diameters. The surface tension of the polymer solution has been shown to be related to the undesirable formation of an array of micron-sized beads present in the electrospun fibers [142,145,147] with the general trend that the thinner the fiber, the smaller the spacing between the beads. Increasing the surface tension tends to make the surface area per unit mass smaller, by changing the shape of the jet into spheres. Concentrated polymer solutions that have yielded fibers by the electrospinning technique have surface tensions in the range of 35–60 dyn=cm. Larrondo and Manley [161] showed that doubling the electric field strength decreased the fiber diameter by a factor of two. However, Baumgarten [158] showed that the diameter of the jet decreased with increasing electric field strength until a minimum value was reached and above which the fiber diameter increased. This effect was caused by the feed rate of the polymer solution through the capillary increasing at high electric fields, which is one of the complexities of the electrospinning process that must be optimized to yield small diameter fibers.
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The majority of research on the production of electrospun fibers has focused on the production of nonwoven mats of random fiber orientation. In order to overcome the random deposition of fibers on the grounded metallic target, the static target shown in Figure 2.32 needs to be replaced by a target that is capable of introducing orientation into the electrospun nanofiber as it comes in contact with the target. In 1982, Bornat [162] described an electrospinning apparatus in which the static metallic target was replaced by a metal cylinder that rotated at 300 rpm. This approach was not specifically aimed at introducing orientation into the electrospun fibers, but to ensure a uniform coating onto a tubular on a vascular prosthetic. Using this approach of a rotating metal cylinder as the target electrode, Kim and Reneker [163] noticed that although the fibers appeared to be randomly oriented in the nonwoven mat, the strength of the material indicated that more fibers lay in the rotating direction. A technique for making continuous uniaxial fiber bundle yarns from electrospun fibers has been described by Smit et al. [164], which consists of spinning onto a water-reservoir collector and drawing the resulting nonwoven web of fibers across the water before collecting the resulting yarn.
2.7.2
Electrospun Polyaniline Fibers
Reneker and Chun [165] first indicated that homogenous polyaniline fibers could be prepared via electrospinning by first dissolving polyaniline in 98% sulfuric acid and then collecting the electrospun fibers using a water coagulation bath. This process was based on the concentrated polyaniline solutions that were wet-spun into textile fibers by Andreatta et al. [54]. This process was later refined by MacDiarmid et al. [166] by a method in which homogenous polyaniline electrospun fibers were prepared from a 20 wt% solution of polyaniline (Versicon from Allied Signal) dissolved in 98% sulfuric acid. The polyaniline nanofibers were prepared with the polymer solution contained in a glass pipette placed 3 cm above the surface of a copper cathode immersed in the water coagulation bath. Using a potential difference between the pipette and the copper cathode of 5 kV, polyaniline nanofibers with an average diameter of 139 nm were formed in the water coagulation bath. The conductivity of these polyaniline nanofibers was found to be 0.1 S=cm. This is lower than the conductivity of the wet-spun polyaniline fibers obtained by Andreatta et al. [54] processed from sulfuric acid (20–60 S=cm) due to partial dedoping of the electrospun fiber arising from the use of a water coagulation bath. Recent work in the acid processing of polyaniline textile fibers by Monkman and coworkers [60,61] has increased the conductivity of the resulting fiber between 100 and 1000 S=cm. These textile fibers were produced from a concentrated polyaniline solution, in which the emeraldine base form of polyaniline was doped with AMPSA and dissolved in DCAA. We have shown in our laboratory that electrically conductive AMPSA-doped polyaniline electrospun fibers can be prepared via electrospinning technique [167]. These electrospun fibers were prepared from an 8 wt% solution of AMPSA-doped polyaniline dissolved in DCAA. The polyaniline nanofibers were prepared with the polymer solution contained in a hypodermic syringe placed 3 cm above the surface of an aluminum foil cathode immersed in a 2-methyl-5-pentanone coagulation bath. Figure 2.33 shows the SEM image of a monofilament of an electrospun AMPSA-doped polyaniline fiber collected in a 2-methyl-5-pentanone coagulation bath at an electric field strength of 2 kV=cm. The diameter of this electrospun polyaniline fiber (1.4 mm) is two orders of magnitude smaller than fibers processed from the same concentrated polyaniline solution using conventional fiber-spinning techniques. A detailed investigation of the temperature dependence of the electrical properties of AMPSA-doped polyaniline electrospun fibers was subsequently performed by Pinto et al. [168]. The polyaniline electrospun fibers were prepared from a 15 wt% solution of AMPSA-doped polyaniline dissolved in DCAA, and were electrospun into an acetone coagulation bath containing an aluminum foil cathode at an electric field strength of 9.3 kV=cm. The conductivity of these polyaniline nanofibers was found to be approximately 4 S=cm at 300 K. After showing a slight initial increase in the conductivity down to 290 K, the conductivity fell by over four orders of magnitude as the temperature was lowered to 10 K. This is contrast to that of a thermally annealed AMPSA-doped polyaniline film, in which the conductivity showed metallic behavior down to about 90 K, below which the conductivity began to drop.
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SEM micrograph of an AMPSA-doped polyaniline electrospun fiber.
Pinto et al. [168] suggested that this difference in the temperature dependence of the electrical properties could be attributed to the rapid evaporation of the solvent during the electrospinning process, resulting in an amorphous rather than a crystalline structure in the fiber as could have been the case in the cast film, where the solvent evaporated slowly leading to larger crystalline regions. Alternately, it could be a result of physical confinement of the polymer chains within the fiber together with dense chain packing due to the electric field used in the electrospinning process that reduces ring flipping that is responsible for charge delocalization. Fabricating homogenous polyaniline electrospun fibers has so far required the use of a coagulation bath due to the low volatility of the solvents used to prepare highly viscous polyaniline solutions. However, nonwoven mats of polyaniline nanofibers have been prepared by blending polyaniline with other polymers that have been successfully electrospun into nonwoven nanofiber mats. Blending of polyaniline with traditional fiber-forming polymers is a commonly used approach to overcome some polyanilines-processing limitations to produce electrically conductive fibers. The initial approach developed by Norris et al. [169] was to blend CSA-doped polyaniline with poly(ethylene oxide) in chloroform. Without the addition of poly(ethylene oxide), no fiber formation occurred because the viscosity and surface tension of the solution were not high enough to maintain a stable drop at the end of the capillary tip. Increasing the concentration of doped polyaniline was found to be insufficient in raising the viscosity and surface tension of the polymer solution to achieve the formation of a stable jet. This is primarily due to the low solubility of CSA-doped polyaniline in chloroform. The diameter of these doped polyaniline=poly(ethylene oxide) electrospun fibers in the nonwoven mat was between 900 nm and 1.9 mm, with a generally uniform diameter along the fiber. It was further found that the UV–vis spectra of these polyaniline electrospun fibers were similar to those for cast films produced from the same solution. It was later shown by MacDiarmid et al. [166] that these doped polyaniline–poly(ethylene oxide) electrospun fibers could be deposited onto an oxidized silicon wafer. They demonstrated that it was possible to measure the conductivity of individual electrospun fibers by depositing gold electrical contacts onto the wafer using a shadow mask evaporation process. It was found that the room temperature conductivity for a 1.3 mm diameter CSA-doped polyaniline–poly(ethylene oxide) electrospun fiber was 33 S=cm, which is two orders of magnitude higher than the conductivity of the corresponding thermally annealed film prepared from the same solution. This result indicates that
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the high electrical fields used to generate the electrospun fibers lead to significant chain alignment and thus give rise to increased electrical conductivity. Subsequent work by Kahol and Pinto [170,171] measured the electron paramagnetic resonance (EPR) of CSA-doped polyaniline–poly(ethylene oxide) (72:28 by weight) electrospun fibers and cast films to determine the mesoscopic structural disorder to explain the increased room temperature conductivity of the electrospun fibers. This study found that the polyaniline blend electrospun fibers exhibited increased Pauli susceptibility, more Lorentzian character in the EPR line shape, and smaller EPR line width. These changes were interpreted as they were associated with increased chain alignment in the electrospun fibers compared with the cast films. Subsequent work has generally focused on decreasing the diameter of these polyaniline blend electrospun fibers. This is because it is relatively easy to prepare electrically conductive blends of CSAdoped polyaniline with a wide variety of conventional polymers. Besides blending doped polyaniline with poly(ethylene oxide), other polymers that have successfully produced electrospun polyaniline fibers include polystyrene, polyacrylonitrile, and poly(methyl methacrylate). For example, Figure 2.34 shows the SEM image of the nonwoven mat of CSA-doped polyaniline=polyacrylonitrile (20:80 by weight) electrospun fibers obtained using an electric field strength of 1 kV=cm. The average diameter of these electrospun polyaniline blend fibers is approximately 200 nm. Unlike the CSA-doped polyaniline– poly(ethylene oxide) blend electrospun fibers, the nonwoven mat of CSA-doped polyaniline–polyacrylonitrile electrospun fibers can be immersed in aqueous solutions without the fibrous network in the nonwoven mat breaking apart due to the dissolution of the poly(ethylene oxide) matrix. Consequently, it is possible to convert the polyaniline–polyacrylonitrile electrospun fibers from their doped state into their insulating emeraldine base oxidation state by immersing the nonwoven mat into a basic aqueous solution (e.g., 0.1 M ammonium hydroxide). The dedoped polyaniline electrospun fibers could then be redoped by immersing the fibers in an acidic solution. Due to the small diameter of these electrospun fibers, and the porous nature of the nonwoven mat, the rate of dedoping and redoping is at least 1–2 orders of magnitude faster than the corresponding cast film of the same thickness. Although the majority of these electrospun polyaniline fibers have submicron dimensions, they cannot be classified as true nanofibers, which is defined as fibers having a diameter less than 100 nm. MacDiarmid et al. [166] first demonstrated that it is possible to fabricate electrospun polyaniline fibers
FIGURE 2.34 SEM micrograph of the electrospun fibers produced from a blend of CSA-doped polyaniline and polyacrylonitrile.
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with diameters less than 100 nm by blending CSA-doped polyaniline with polystyrene. Electrospinning a solution of CSA-doped polyaniline–poly(ethylene oxide) (20:80 by weight) dissolved in chloroform resulted in an average diameter of 86 nm for the resulting electrospun fibers, with the largest fiber having a diameter of 100 nm. It was shown in subsequent work that is possible to prepare CSA-doped polyaniline–poly(ethylene oxide) nanofibers with diameters between 5 and 70 nm [172]. For comparison, it should be noted that the diameters of multiwalled carbon nanotubes are in the range of 4–30 nm. These electrospun polyaniline nanofibers were obtained by decreasing the concentration of poly(ethylene oxide) dissolved blend to 4% by weight. The CSA-doped polyaniline–poly(ethylene oxide) nanofibers were spun from a chloroform solution using an electric field strength of 0.26 kV=cm, with the electrospun nanofibers collected onto an oxidized silicon wafer. Scanning conductance microscopy of these electrospun nanofibers shows that fibers with diameter below 15 nm are electrically insulating. It was proposed that the small diameter might enable complete dedoping in air, or may be smaller than phase-separated grains of doped polyaniline and poly(ethylene oxide). Electrical contacts to nanofibers enabled single fiber I–V characteristics, which showed that the conductivity of the electrospun nanofibers decreased as the diameter of the fiber decreased. Measurements of I–V characteristics of single nanofibers with sharply varying diameter were found to be rectifying, which is consistent with formation of Schottky barriers at the nanofiber–metal contacts. The final approach that has been used to fabricate electrospun fibers is based on replicating the early methods of fabricating composite polyaniline textile fibers. This work involved depositing a thin layer of doped polyaniline onto a textile fiber using either an in situ polymerization technique or immersing the textile fiber into a solution of polyaniline dissolved in an organic solvent. Dong et al. [173] demonstrated that it was possible to fabricate composite HCl-doped polyaniline nanofibers by immersing a nonwoven mat of 290 nm diameter poly(methyl methacrylate) electrospun fibers in an aqueous solution containing aniline, hydrochloric acid, and ammonium persulfate. After a polymerization time of 8 min, a 30 nm thick polyaniline layer was deposited onto the electrospun poly(methyl methacrylate) fibers and the nonwoven fiber mat had a conductivity of 0.3 S=cm. An interesting derivative of this work is a process in which polyaniline nanotubes were fabricated by Dong et al. [174] by first coating an electrospun fiber with a thin polyaniline layer using an in situ polymerization technique followed by thermolytic removal of the core polymer to form the polyaniline nanotube. In this work, poly(L-lactide) was selected as the fiber template, since it can be readily processed into fibers with diameters in the submicrometer range by electrospinning and has a relatively low decomposition temperature (2358C–2558C). This low decomposition temperature reduces the possibility of cross-linking or other structural damage of the polyaniline backbone that would reduce the conductivity of the polyaniline nanotube during thermolysis of the core. Varying the solution stoichiochemistry and the deposition time can control the wall thickness of the polyaniline layer deposited onto the electrospun fiber. After the thermal processing step, it was necessary to redope the polyaniline nanotubes using an acidic solution since the elevated temperatures also removed the dopant anion from the polyaniline backbone. The conductivity of the composite polyaniline electrospun fibers was 0.4 S=cm, and the conductivity of the polyaniline nanotubes was 0.3 S=cm. Choi et al. [175] showed that composite polyanline nanofibers could be fabricated in which doped polyaniline is deposited onto electrospun silica nanofibers using either an in situ polymerization technique, or immersing the electrospun nonwoven mat into a solution of doped polyaniline dissolved in chloroform or NMP. The silica nanofibers were electrospun from a ripened tetraethyl orthosilicate solution. It was found that polyaniline composite electrospun fibers prepared by the in situ polymerization had a more uniform polyaniline coating than those prepared by immersing the nonwoven mat into the polyaniline solution. 2.7.2.1 Electrospun Polypyrrole Fibers Although we have been able to electrospin polyaniline into nanofibers, polypyrrole is unable to be processed into homogenous textile fibers since this ICP is insoluble in common organic solvents and infusible, thus making thermal or solution processing difficult. To overcome this limitation,
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FIGURE 2.35
SEM micrograph of polypyrrole-coated electrospun polyacrylonitrile nanofibers.
MacDiarmid et al. [166] deposited 25 nm thick layer of chloride-doped polypyrrole onto polyacrylonitrile electrospun fibers through an in situ polymerization technique by immersing an electrospun nonwoven mat in a solution containing the pyrrole monomer, an oxidant, and the desired dopant anion. In this initial work, chloride derived from the ferric chloride oxidant was used as the dopant anion. Similar to what has been observed for polypyrrole-coated textile fibers, replacing the chloride dopant anions with aromatic sulfonated ions to dope the polypyrrole backbone not only lowers the surface resistivity of the polypyrrole composite electrospun nanofibers by two orders of magnitude (106 to 104 V=x) but also improves their thermal stability. For example, Figure 2.35 shows the morphology of a nonwoven mat consisting of polyacrylonitrile electrospun fibers with an average diameter of 500 nm coated with a 30 nm thick layer of p-toluenesulfonate-doped polypyrrole. It was later shown by Kang et al. [176] that homogenous polypyrrole electrospun fibers could be prepared by controlling the chemical synthesis process to form a low-molecular-weight polypyrrole doped with dodecylbenzenesulfonate, which is soluble in a range of organic solvents. The homogenous polypyrrole electrospun fibers were then prepared from a 35 wt% solution of dodecylbenzenesulfonatedoped polypyrrole dissolved in chloroform with an electric potential of 30–45 kV applied between the polymer solution contained in a hypodermic syringe and a rotating drum cathode. After removing the excess dodecylbenzenesulfonic acid with methanol, the polypyrrole electrospun fibers possessed an average diameter of 3 mm with an electrical conductivity of the nonwoven mat of 0.5 S=cm. The conductivity of the nonwoven mat was greater than that of the powder and cast film (0.2–0.3 S=cm), which may indicate that molecular orientation of the polymer chains was induced during the electrospinning process. The quality of the nonwoven mat was greatly improved in terms of reduced fiber diameter and fiber uniformity by the addition of 20% of poly(vinyl cinnamate) by weight to increase the viscosity of the concentrated polypyrrole solution.
2.7.3
Electrospun MEH–PPV Fibers
One of the other ICPs that has received considerable attention for processing into electrospun fibers is poly(2-methoxy-5-(20 -ethylhexyloxy)-1,4-phenylene vinylene) (MEH–PPV) because of its ability to form concentrated solutions (>5 wt%) in organic solvents such as chloroform, 1,2-dichloroethane,
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benzene, toluene, and tetrahydrofuran. There has been much interest in this soluble poly(p-phenylene vinylene) derivative due to its potential use in LEDs and photovoltaic devices. Madhugiri et al. [177] reported that homogenous electrospun MEH–PPV fibers could be spun from a 7–10 wt% solution in 1,2-dichloroethane at an electric field strength of 0.9 kV=cm, since the boiling point of 1,2-dichloroethane (768C–788C) is the most conductive of the above solvents for electrospinning. The resulting MEH–PPV electrospun fibers were found to be approximately 200 nm in diameter with many of the fibers exhibiting a leaflike structure, which appeared when two fibers are fused together. The low viscosity of the MEH–PPV electrospinning solution (380–400 cP), which is lower than the typical range of 1000–2000 cP, appears to be the cause of these leaflike structures. It should be noted that the viscosity could not be further increased since 10 wt% is the upper solubility limit of MEH–PPV in 1,2-dichloroethane. To overcome this limitation, Madhugiri et al. [177] employed a novel dual syringe electrospinning technique to produce composite MEH–PPV=mesoporous silica fibers. The dual syringe electrospinning technique involves placing the two syringe barrels, one for each solution, with one on the top of the other such that the flat tips of the hypodermic needles are touching each other. The diameter of these composite MEH–PPV=mesoporous silica electrospun fibers was between 700 and 800 nm, but the amount of leaflike structures that was observed in the homogenous MEH– PPV electrospun fibers was markedly reduced. Furthermore, it was found in photoluminescence studies that the composite MEH–PPV electrospun fibers exhibited a blue shift in the polymer emission as compared to pure MEH–PPV electrospun fibers (535–566 nm) when excited at 550 nm. The blue shift observed in the case of composite electrospun fibers could be attributed to a disruption in the conjugation length, or prevention of aggregation in the polymer when electrospun along with the mesoporous silica since both of these reasons are known to cause a blue shift in the emission of fluorescent polymers. Later it was shown by Wutticharoenmongkol et al. [178] that MEH–PPVelectrospun fibers with average diameters ranging from 300 nm to 5.1 mm could be prepared from solutions of MEH–PPV blended with polystyrene when dissolved in chloroform, 1,2-dichloroethane, and tetrahydrofuran using electric field strength of 1.5 kV=cm. The photoluminescence spectra of the MEH–PPV electrospun fibers exhibited a slight red shift in the emission when compared to the corresponding solutions (560–570 nm). One of the most effective approaches to electrospinning ICPs into nanofibers has been to blend the ICP with other electrospinnable fibers in solution. Li et al. [179] recently showed that this is not only the approach to fabricate electrospun MEH–PPV fibers. It was found that defect-free MEH–PPV electrospun fibers could not be fabricated by codissolving MEH–PPV and poly(vinyl pyrrolidinone) in chloroform because chloroform is a poor solvent for electrospinning poly(vinyl pyrrolidinone) and the solvents most appropriate for producing poly(vinyl pyrrolidinone) electrospun fibers such as ethanol–water mixtures, caused the MEH–PPV to precipitate. To overcome these limitations, it was shown that composite MEH–PPV electrospun fibers could be prepared using an electrospinning spinneret comprised of two coaxial capillaries. The spinneret is fabricated by inserting a polymer-coated silica capillary into a stainless steel needle, and is commonly used to prepare core-sheath and hollow fibers. In a typical procedure, an MEH–PPV solution in chloroform and a poly(vinyl pyrrolidinone) solution in the ethanol–water mixture were simultaneously fed through the inner and outer capillaries, respectively. The diameter of these electrospun fibers was found to be between 100 and 500 nm. Fluorescence optical microscopy indicated that MEH–PPV had been incorporated homogenously into the poly(vinyl pyrrolidinone) electrospun fiber. After the poly(vinyl pyrrolidinone) has been removed by ethanol extraction, electrospun fibers with homogenous fluorescence emission characteristic of MEH– PPV remained. The photoluminescence emission spectra of the electrospun MEH–PPV fibers is similar to that of a spin-coated MEH–PPV thin film from chloroform whereas the UV–vis absorption spectra indicated that the absorption peak is slightly red-shifted compared to a spin-coated MEH–PPV film from chloroform. This red shift implies that the MEH–PPV chains in the electrospun fiber possess a more extended chain conformation and more delocalized p-conjugation. This increase in conjugation length has been ascribed to significant stretching of the MEH–PPV chains in the liquid jet during the electrospinning process.
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2.7.4
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Devices Fabricated from Conducting Polymer Electrospun Fibers
The majority of research has focused on the fabrication of ICP electrospun fibers but potential applications of these materials have been gradually appearing. These applications generally exploit either the high surface that can be obtained using electrospun fibers or the unique 1D geometry that can be obtained using electrospun fibers. One of the most important polymeric devices fabricated from ICP has been the FET, since it forms the basic building block for the fabrication of logic circuits and switches for displays. Traditionally, FET devices are fabricated using a planar 2D ICP layer deposited onto a prepatterned silicon wafer. As noted earlier, ICP electrospun fibers have been deposited onto prepatterned oxidized silicon wafers to measure the conductivity of an individual electrospun fiber. The same approach has been used to prepare ICP nanofiber FET devices. Compared to the traditional 2D FET devices, the electrospun nanofiber FET devices have a 1D geometry similar to devices made from carbon nanotubes. Pinto et al. [180] reported the first electrospun FET device behavior, in which electrospun CSA-doped polyaniline–poly(ethylene oxide) nanofibers layer were deposited onto the prepatterned silicon wafer. Saturation channel currents were observed at surprisingly low source drain voltages. Since the poly(ethylene oxide) component of the electrospun fiber is insulating, it is anticipated that higher mobility will be observed by reducing or eliminating poly(ethylene oxide) from the fiber that decreases the barriers to charge transport between polyaniline chains. Babel et al. [181] later reported the performance of FET devices from electrospun fibers prepared from a binary blend of MEH–PPV with regioregular poly(3-hexylthiophene). The electrospun fibers that were deposited onto the prepatterned silicon wafer were prepared using the previously described two coaxial capillary spinneret in which the blend solution of MEH–PPV and regioregular poly(3-hexylthiophene) dissolved in chloroform and the poly(vinyl pyrrolidinone) solution in the ethanol–water mixture were simultaneously fed through the inner and outer capillaries, respectively. Poly(vinyl pyrrolidinone) was subsequently removed by immersing the silicon wafer in ethanol to form electrospun blend fibers of MEH–PPV and regioregular poly(3-hexylthiophene). FET devices fabricated from nonwoven mats of these electrospun nanofibers showed p-channel transistor characteristics with hole mobility in the range of (0.05–1) 104 cm2=Vs depending on the concentration of MEH–PPV in the electrospun fiber. If corrected for the reduced channel area of the transistors, the effective hole mobility in these blends are actually one order of magnitude higher (0.05–1) 103 cm2=Vs, since the nonwoven mats of electrospun nanofibers occupied only 10% of the channel area. These effective hole mobility values are similar to that obtained using spin-coated blend thin films. It has been recently shown that the sensitivity and response time of ICP fiber chemi-resistive sensors are governed by the morphology of the fiber detector, with faster response times obtained when the detector possesses a high surface-to-volume ratio through the use of ICP electrospun nanofibers [169,182]. For example, it has been demonstrated that the detection of ammonia vapor was an order of magnitude faster when using a nonwoven mat of CSA-doped polyaniline–poly(ethylene oxide) electrospun nanofibers than a thin film of doped polyaniline with the same composition (few seconds vs. 1–2 min). This presumably arises from a faster diffusion rate of ammonia vapor into the polyaniline nanofibers. The above condition highlights the enormous effect on increase in the surface-to-volume ratio accomplished by the electrospun fibers can have on selected chemical properties of a specific polymer. Similar improvements in sensitivity and response time have been reported for chemically synthesized polyaniline nanofibers for the detection of ammonia and HCl vapors [183]. It has been envisaged that the next generation of nonwoven textiles for chemical protection could embed these ICP nanofibers into the fabric to both simultaneously detect and neutralize harmful vapors [184]. Similarly, Aussawasathien et al. [185] compared the performance of CSA-doped polyaniline–polystyrene electrospun nanofibers with a thin film of the same composition for the electrochemical detection of hydrogen peroxide. As expected, the thin-film sensor showed significantly weaker currents than those of the electrospun nanofibers, with both materials showing a linear response of the redox current as a function of hydrogen peroxide concentration. However, the electrospun nanofiber sensor showed a much higher sensitivity, as evidenced by the greater slope. It was also shown that the detection of
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glucose could be achieved using glucose oxidase-immobilized polyaniline electrospun nanofibers. The amperometric response from the glucose oxidase-immobilized polyaniline electrospun nanofiber sensor was higher than that of the traditional film sensor fabricated from the same materials. Since the nanofibers possess a higher surface area, the amount of glucose oxidase that could be immobilized at the sensor surface is higher than that for the film sensor and this is responsible for the increased sensitivity of the nanofiber sensor. The well-known fact that the rate of electrochemical reactions is proportional to the surface area of the electrode makes ICP electrospun fibers ideal candidates as electrode materials in other small electrochemical devices such as supercapacitors and batteries. The small diameter of the fibers makes it possible for ions to rapidly diffuse between the center of the fiber and the surrounding electrolyte, which should lead to enhanced performance of electrochemical devices constructed from these electrospun fibers. Although no energy storage devices have so far been fabricated that utilize ICP electrospun fiber electrodes, there have been reports of using graphitized polyacrylonitrile electrospun fibers as the electrode material for fabricating carbon-based supercapacitors [186]. Recently, Li et al. [187] electrospun polyaniline blended with a natural protein, gelatin into nanofibers, with the goal of investigating their use as a conductive scaffold for tissue engineering purposes. To test the usefulness of PANI–gelatin blends as a fibrous matrix for supporting cell growth, H9c2 rat cardiac myoblast cells were cultured on fiber-coated glass-cover slips. Their results indicated that the PANI–gelatin blend electrospun fibers supported H9c2 cell attachment and proliferation to a similar degree as the control tissue culture-treated plastic and smooth glass substrates. Depending on the concentrations of PANI in the electrospun fibers, the cells initially displayed different morphologies on the fibrous substrates, but after 1 week all cultures reached confluence of similar densities and morphology. Taken together, these results suggest that PANI–gelatin blend nanofibers might provide a novel conductive material well suited as biocompatible scaffolds for tissue engineering.
2.8
Concluding Remarks
Many of the issues associated with the processing of polyaniline fibers are only now addressed based on pilot scale trials. Many of these issues arise from the different lengths of time during which the fiber dope solution must possess stable rheological properties to produce a fiber with consistent mechanical and electrical properties. This length of time can vary from a few hours required for academic studies to several days for pilot plant production runs. While the majority of the work that has been published has been largely based on processing the polymer in its neutral state (emeraldine base or leucoemeraldine) since the polymer is not a rigid rod, these fibers generally suffer from poor mechanical properties due to the formation of voids when produced on a pilot plant scale. Research is needed to reduce or eliminate the formation of voids in these fibers by controlling the fiber formation kinetics in the coagulation bath. This should lead to the formation of base-processed fibers with improved mechanical and electrical properties. Additionally, one of the major limitations of the base-processed polyaniline fibers is that the highest conductivity that has been obtained from a base-processed polyaniline fiber after drawing is 350 S=cm [36]. This is significantly lower than which has been achieved for stretched polyaniline films and fibers obtained using an acid-processing route (400–1000 S=cm). The acid-processing route for spinning polyaniline into fibers has proven to be the catalyst for developing the next generation of polyaniline fibers with a superior combination of both mechanical and electrical properties. Based on this foundation, polyaniline fibers and hollow fibers have now been prepared on a pilot plant scale and it has been shown that these materials can be readily processed into electrically conductive textiles by weaving, knitting, stitching, and braiding using conventional textile-processing equipment. In contrast, the in situ polymerization approach for depositing an intractable ICP, such as polypyrrole, onto an existing textile substrate has still been the most effective approach for fabricating electrically conductive textiles from these intractable ICPs. This technique has been successfully converted from laboratory into commercial production for producing large quantities of these textiles using a relatively inexpensive aqueous-based process.
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Processing improvements have enabled the pilot plant to commercial scale production of ICP-based electrically conductive fibers and fabrics and it is likely that these materials will become commercially important in the emerging field of ‘‘smart fabrics and intelligent textiles.’’ The current stage of development has permitted the fabrication of prototype devices in which active functions have been successfully embedded into the garment (e.g., resistive heating, vapor detection, strain and force sensing). As the stability concerns of these ICP-based electrically conductive textiles are addressed further and the economy of scale lowers the manufacturing cost, significant market penetration devices should take place.
Acknowledgments We would like to thank the many current and former staff at Santa Fe Science Technology for their contributions to the work that has been described in this chapter. They include Drs. Phil Adams, Danila Bowman, Lori Brown, Andrei Fadeev, Russell Goering, Wen Lu, John Pellegrino, Baohua Qi, Elisabeth Smela, Dali Yang, and Guido Zuccarello. The majority of the work described in this chapter was supported by the Defense Science Office of the Defense Advanced Research Projects Agency for which the authors are extremely grateful.
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160. Buchko, C.J., K.M. Kozloff, and D.C. Martin. 2001. Surface characterization of porous, biocompatible protein polymer thin films. Biomaterials 22:1289. 161. Larrondo, L., and R.S.J. Manley. 1981. Electrostatic fiber spinning from polymer melts. I. Experimental observations on fiber formation and properties. J Polym Sci Polym Phys 19:909. 162. Bornat, A. 1982. Electrostatic spinning of tubular products. US Patent 4,323,525. 163. Kim, J.-S., and D.H. Reneker. 1999. Polybenzimidazole nanofiber produced by electrospinning, Polym Eng Sci 39:849. 164. Smit, E., U. Buttner, and R.D. Sanderson. 2005. Continuous yarns from electrospun fibers. Polymer 46:2419. 165. Reneker, D.H., and I. Chun. 1996. Nanometre diameter fibres of polymer, produced by electrospinning. Nanotechnology 7:216. 166. MacDiarmid, A.G. Jr., W.E. Jones, I.D. Norris, J. Gao Jr., A.T. Johnson, N.J. Pinto, H. Hone, B. Han, F.K. Ko, H. Okuzaki, and M. Llaguno. 2001. Electrostatically-generated nanofibers of electronic polymers. Synth Met 119:27. 167. Norris, I.D., and B.R. Mattes. 2006. In preparation for submission to Synthetic Metals. 168. Pinto, N.J., P. Carrio´n, and J.X. Quin˜ones. 2004. Electroless deposition of nickel on electrospun fibers of 2-acrylamido-2-methyl-1-propanesulfonic acid doped polyaniline. Mater Sci Eng A 366:1. 169. Norris, I.D., M.M. Shaker, F.K. Ko, and A.G. MacDiarmid. 2000. Electrostatic fabrication of ultrafine conducting fibers: Polyaniline=polyethylene oxide blends. Synth Met 114:109. 170. Kahol, P.K., and N.J. Pinto. 2002. Electron paramagnetic resonance investigations of electrospun polyaniline fibers. Solid State Commun 124:195. 171. Kahol, P.K., and N.J. Pinto. 2004. An EPR investigation of electrospun polyaniline–polyethylene oxide blends. Synth Met 140:269. 172. Zhou, Y., M. Freitag, J. Hone, C.A.T. Staii, J. Johnson, N.J. Pinto, and A.G. MacDiarmid. 2003. Fabrication and electrical characterization of polyaniline-based nanofibers with diameter below 30 nm. Appl Phys Lett 83:3800. 173. Dong, H., V. Nyame, A.G. MacDiarmid, and W.E. Jones Jr. 2004. Polyaniline=poly(methyl methacrylate) coaxial fibers: The fabrication and effects of the solution properties on the morphology of electrospun core fibers. J Polym Sci 42:3934. 174. Dong, H., S. Prasad, V. Nyame, and W.E. Jones Jr. 2004. Sub-micrometer conducting polyaniline tubes prepared from polymer fiber templates. Chem Mater 16:371. 175. Choi, S.-S., B.Y. Chu, D.S. Hwang, S.G. Lee, W.H. Park, and J.K. Park. 2005. Preparation and characterization of polyaniline nanofiber webs by template reaction with electrospun silica nanofibers. Thin Solid Films 477:233. 176. Kang, T.S., S.W. Lee, J. Joo, and J.Y. Lee. 2005. Electrically conducting polypyrrole fibers spun by electrospinning. Synth Met 153:61. 177. Madhugiri, S., A. Dalton, J. Gutierrez, J.P. Ferraris, and K.J. Balkus Jr. 2003. Electrospun MEHPPV=SBA-15 composite nanofibers using a dual syringe method. J Am Chem Soc 125:14531. 178. Wutticharoenmongkol, P., P. Supaphol, T. Srikhirin, T. Kerdcharoen, and T. Osotchan. 2005. Electrospinning of polystyrene=poly(2-methoxy-5-(2-ethylhexyloxy)-1,4-phenylene vinylene) blends. J Polym Sci 43:1881. 179. Li, D., A. Babel, S.A. Jenekhe, and Y. Xia. 2004. Nanofibers of conjugated polymers prepared by electrospinning with a two-capillary spinneret. Adv Mater 16:2062. 180. Pinto, N.J., A.T. Johnson Jr., A.G. MacDiarmid, C.H. Mueller, N. Theofylaktos, D.C. Robinson, and F.A. Miranda. 2003. Electrospun polyaniline=polyethylene oxide nanofiber field-effect transistor. Appl Phys Lett 83:4244. 181. Babel, A., D. Li, Y. Xia, and S.A. Jenekhe. 2005. Electrospun nanofibers of blends of conjugated polymers: Morphology, optical properties, and field-effect transistors. Macromolecules 38:4705.
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182. MacDiarmid, A.G., I.D. Norris, W.E. Jones, Jr., M.A. El-Sherif, J. Yuan, B. Han, and F.K. Ko. 2000. Polyaniline based chemical transducers with sub-micron dimensions. Polym Mater Sci Eng 83:544. 183. Huang, J., S. Virji, B.H. Weiller, and R.B. Kaner. 2003. Polyaniline nanofibers: Facile synthesis and chemical sensors. J Am Chem Soc 125:314. 184. Adam, D. 2001. A fine set of threads. Nature 411:236. 185. Aussawasathien, D., J.-H. Dong, and L. Dai. 2005. Electrospun polymer nanofiber sensors. Synth Met 154:37. 186. Kim, C., and K.S. Yang. 2003. Electrochemical properties of carbon nanofiber web as an electrode for supercapacitor prepared by electrospinning. Appl Phys Lett 83:1216. 187. Li, M., Y. Guo, Y. Wei, A.G. MacDiarmid, and P.I. Lelkes. 2006. Electrospinning polyanilinecontained gelatin nanofibers for tissue engineering applications. Biomaterials 27:2705.
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3 Inkjet Printing and Patterning of PEDOT–PSS: Application to Optoelectronic Devices
Yuka Yoshioka and Ghassan E. Jabbour
3.1
3.1 3.2 3.3 3.4
Introduction......................................................................... 3-1 Direct Printing of PEDOT–PSS Layers ............................. 3-5 Deactivatived Patterning of PEDOT–PSS ....................... 3-10 Conclusion ......................................................................... 3-20
Introduction
The development of electronics based on conducting polymers has gained more attention due to their light-weight, mechanical flexibility, and simple processing as well as their optoelectronic properties. An advantage of these polymers is that their optoelectronic properties can be modified by designing the chemical function on the molecules, the alignment of polymer chains, and doping conditions. Poly(3,4-ethylenedioxythiophene)–poly(styrene sulfonate) (PEDOT–PSS, Figure 3.1) was first synthesized in the late 1980s, and is one of the most successful conducting polymers to be developed and studied. PEDOT is prepared by standard oxidative chemical or electrochemical polymerization methods. PEDOT, itself, is found to be highly conductive (400–600 S cm1), highly transparent, electrochemically stable, and thermally stable (up to 2308C) in thin, oxidized films (doped with PF6, BF4, or CF3SO3). However, the processability of PEDOT is very poor because it is an insoluble polymer. This drawback can be overcome by polymerizing it in combination with a water-soluble polyelectrolyte, PSS. The resulting PEDOT–PSS is dark blue, electrochemically stable in its p-doped form, moderately transparent with high electrical conductivity (1–10 S cm1) [1], and has excellent film-forming properties. Since conjugated polymers are generally polycrystalline materials, the bulk conductivity is defined by charge transport properties along the polymer chain, from one chain to another chain and across domain boundaries. Kirchmeyer and Reuter suggested that PEDOT is mainly composed of 6–18 repeating units, which were tightly attached to the PSS chains of higher molecular weight (Figure 3.2) [2]. This usually creates a strong dependence between the conductivity and the morphology of the 3-1
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O
O
O
O
O
S
n
S O
O
S
S
FIGURE 3.1
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O
S O
O
n
SO3
The molecular structure of poly(3,4-ethylenedioxythiophene)–poly(styrene sulfonate) (PEDOT–PSS).
polymer film. Recent studies have shown that the conductivity of PEDOT–PSS films is considerably increased by adding high boiling point liquid additives, such as glycerol [3–6], sorbitol [7,8], N-methylpyrrolidone (NMP) [9], ethylene glycol, or meso-erythriol [10], during the solution processing, followed by a heat treatment (Figure 3.3). This improvement is due to the morphological changes of the PEDOT– PSS (relaxation of the polymer chains) induced by the addition of the solvent. Therefore, the additives partially act as plasticizers, which lower the glass transition temperature (Tg) to form a more homogeneous distribution between conducting PEDOT chains during the drying process. Thus, the film morphology and chemical and physical structures of the conducting polymers can be strongly influenced by a variety of postdeposition treatments, such as heat treatments [11–13]. It should be noted that due to the strong ionic bonding interaction between PEDOT and PSS, PEDOT–PSS does not have a specific Tg value (Figure 3.4). However, its electrical property and morphology change little over a wide range of temperatures until it reaches a degradation temperature of 2308C. Furthermore, adding a polar solvent, for instance, dimethyl sulfoxide (DMSO), N,N-dimethyl formamide (DMF), or tetrahydrofuran (THF) increases the conductivity of the resulting films (Figure 3.5) [14]. A possible explanation for this is that the polar solvent partially dissolves the PEDOT in the PEDOT–PSS complex, enhancing a morphological rearrangement and clustering for easier charge transportation [2]. Due to the selective nature of PEDOT–PSS, it has been used in several industrial and laboratory applications in areas from biology to engineering. These include antistatic coatings in photographic films (Agfa), electrochromic displays, organic electrochemical transducers, organic thin film transistors, electrochemical organic transistors, organic RC filter circuits, polymer capacitors, polymer-dispersed liquid-crystal displays, flexible organic speakers, infrared detectors, ion-selective electrodes, a hole injection layer or anode of organic light-emitting devices (OLEDs), and a hole-collecting layer or electrode in organic photovoltaic devices.
PSS unit PEDOT unit
Ionic bonding
FIGURE 3.2
Schematic image of PEDOT–PSS chain: a long PSS chain with PEDOT oligomeric chains.
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CH3
OH OH
N
OH
OH
HO
3-3
O
HO
OH
OH OH
Glycerol b.p. = 290°C
N-methylpyrrolidone b.p. = 202°C
Sorbitol b.p. = 295°C
OH OH
OH
HO
HO HO meso-erythriol b.p. = 330°C
Ethylene glycol b.p. = 198°C
FIGURE 3.3
Molecular structures of glycerol, sorbitol, N-methylpyrrolidone, ethylene glycol, and meso-erythriol.
0
Heat flow (W/g)
−0.1 −0.2 −0.3 −0.4 −0.5 −0.6
0
100
200
300
400
Temperature (⬚C)
FIGURE 3.4 Differential scanning calorimetry (DSC) result for pristine PEDOT–PSS (2920 modulated DSC TA instrument).
O
O
S H3C
O
C CH3
DMSO m = 3.96 D b.p. = 189⬚C
H
N(CH3)2
DMF m = 3.82 D b.p. = 153⬚C
THF m =1.75 D b.p. = 75⬚C
FIGURE 3.5 Molecular structures, dipole moment, and boiling point of DMSO, DMF, and THF. *Reference dipolemoment: m (water) ¼ 1.85 D (Lide, D.R., Handbook of Chemistry and Physics, CRC Press, Boca Raton, FL, 2003–2004).
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Although PEDOT–PSS has a higher sheet resistance compared to the state-of-the-art transparent conductor indium tin oxide (ITO), it is possible to replace ITO by a conducting polymer in some applications. Especially in low information content flexible display applications, ITO is found to develop cracks upon repeated flexing, thus losing conductivity, which can be overcome by using polymer-based materials, due to the benefit of good mechanical flexibility. Also, using ITO requires costly sputtering methods, which could be simplified by using solution processing techniques. Such techniques encompass not only spin-coating, but also printing methods, such as screen printing [15,16], line patterning [17], or inkjet printing. To maximize the resolution, the following methods are often combined with various patterning techniques: UV lithography, soft lithography [18], imprint lithography [19], and deactivation process with the oxidizing agent (i.e., sodium hypochlorite) by screen printing [20] or inkjet printing [21]. As a buffer layer, PEDOT offers the possibility of planarizing the ITO surface, thus reducing leakage currents and increasing fabrication yield. A 50–80 nm thick layer of PEDOT–PSS can provide assistance to both smoothing out the rough ITO surface as well as enhancing hole injection into OLED or PLED structures. Inkjet printing is one of the most promising deposition techniques, which has been used in a variety of fields including ceramics, metals, organic semiconductors, and biopolymers [22–24]. It has the advantage of being fast and simple with high throughput. A desktop computer and a standard inkjet printer can control the delivery of picoliter volumes of liquids in precise patterns. Customized industrial inkjet printers have matured to impressive performances. Nowadays, an industrial inkjet printer is capable of depositing a subarray of polymer droplets within pixels over a large high-resolution flat panel display. Inkjet printing does not consume significant amount of material, as is the case with spin-coating. It promises a low-cost, maskless, and noncontact patterning approach. Inks that are based on PEDOT– PSS have been extensively used to fabricate microelectromechanical and chemical structures and devices with conducting polymer-based electronic parts: OLEDs [25–29], organic thin film transistors [30], allpolymer capacitors [31], polymer-based rectifying diodes [32], electrochromic displays [33], chemical sensing devices (chemical fuse) [34], and glucose biosensor prototypes [35]. In building device structures by inkjet printing, there are a number of approaches, for example, forming individual dots, lines, areas, or thick layers. Generally, by arranging the printing conditions, one can connect individual dots to form a line. By narrowing the offset of line alignment, one can connect lines to form a layer. Furthermore, in office-type inkjet printer, about 50–100 mm diameter dots are dispersed over the area of the printed image and the amount of dot overlap increases as the original density of dots increases. The shape, thickness, and surface morphology of droplets and layers are greatly influenced by energetics of the substrate surface and ink, as well as the process of solvent evaporation. One of the most extensive applications of inkjet printing is pixel printing for OLED applications [27–29]. Many companies including Cambridge Display Technology, Seiko-Epson, and Toshiba demonstrated display prototypes ranging in size from several to 40 in., in which the polymeric color pixels and base (hole injection) layers were deposited. In these displays, inkjet printer was used to dispense polymers in a predefined pixel pattern (30 mm pixel size) on photolithographically patterned substrates. For their applications, the uniformity of inkjet-printed features is of a particular concern [36]. Generally, a circular drop on the solid surface evaporates from the edges and flows outward. This capillary flow tends to carry most of the dispersed materials to the edge, a phenomenon known as the coffee-ring effect [37]. This leads to a nonuniform thickness profile of a dried droplet with a thick ring around the edge of the drop [38]. The capillary flow rate can be modified by controlling the solute concentration and molecular weight, surface tension, viscosity of the liquid, and drying conditions. A higher solute concentration or faster drying process tends to push edges of droplets more outwards [39–41]. However, if polymer–substrate interaction is stronger than interpolymer interactions, it is possible to immobilize the polymer on the substrate after printing [42]. This can be achieved using high-viscosity solvents, which hinders the capillary flow of solutes within a drop as it dries on the substrate. For noncircular droplets, including a line and layer of liquid, deposition behavior can be more complicated. In fact, it has been extremely difficult to achieve uniform line or layer structures on uncoated glass or plastic
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substrates, since the line or the layer formed by joining droplets often increases surface roughness and causes pin-hole formation on the printed surface during the drying process. In order to place drops in precise locations, a surface-energy pattern can be defined on the glass substrates before performing inkjet printing (surface-energy assisted inkjet) [30]. For example, hydrophilic glass substrates can be patterned by polyimide (hydrophobic polymer), using a photolithographic process. Since surface energy is the dominant force for causing spreading liquid, a hydrophobic bank efficiently prevents a drop from spreading beyond the confined areas of hydrophilic surface. Fine feature sizes of several microns to submicron have been successfully demonstrated. In this chapter, we discuss methods of direct patterning of conducting polymers to create a variety of images including gray scale by inkjet printing. Well-established deskjet printing technology allows us to simplify the modification process of printing parameters. We will present two ways of achieving this—one by the direct deposition of PEDOT–PSS and the other by modifying the sheet resistivity of already-deposited PEDOT–PSS electrodes by using oxidizers, such as sodium hypochlorite or hydrogen peroxide. In comparison with abovementioned inkjet printing methods, our approach requires no prepatterning of the substrate (no photolithography). For the design of layer structures, we utilized HSL (hue, saturation, and luminosity) color function on the Power Point software along with office-type thermal inkjet printers.
3.2
Direct Printing of PEDOT–PSS Layers
This study was carried out using a modified Hewlett Packard (HP) thermal desktop inkjet printer. Glass substrates were fixed to the transparency sheet with kapton tape, and the sheets were inserted into the printer. The HP driver software was used to select the printing positions. A pattern was designed using Microsoft Power Point software. A PEDOT–PSS ink was formed, consisting of 86 vol.% of Baytron P, 4.5 vol.% of glycerol, 0.045 vol.% of triton X-100 (surfactant), and 9.5 vol.% of water. Glycerol was added to enhance the conductivity of formed films as well as a humectant to avoid nozzle clogging during the printing process. The thickness of the inkjet printing films was controlled through the HSL function using the existing Power Point values. In this case, the ‘‘saturation’’ (S), which controls the respective quantity of RGB (red, green, and blue) color, was held at 0. The hue (H), which identifies the composition of the color, was also held constant. Only ‘‘luminosity’’ (L), which dictates the darkness of the color itself (and thus the amount of printed PEDOT–PSS ink), was changed over six different values as shown in Figure 3.6. The density of droplets defines the gray scale of image, which depends upon the surface energy of the ink. The PEDOT–PSS ink described above has a low surface energy, and therefore forms a fully interconnected film.
Luminosity (L) 0 10 20 30 50 70
(a) Gray values
FIGURE 3.6 not defined.
(b) Analogy of printing dots on paper
(a) Examples of the gray scale and (b) HSL values used to program ink loads. S ¼ 0 and H was
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Thickness (nm)
200
150
100
50
0 0
10
20
30
40
60
50
Luminosity
FIGURE 3.7 Measured film thicknesses for various luminosity values. L ¼ 0–50. Film became discontinuous for L 70 (72% darkness).
Continuous conductive films were formed with similar thicknesses since there is only one ink. The continuity of films was visually inspected after the PEDOT–PSS ink was printed. After printing, substrates were immediately placed on a hot plate at 1108C for 2 h in ambient atmosphere to dry out residues of glycerol and water. Film thickness, sheet resistivity, and surface roughness average were then measured. As shown in Figure 3.7, a systematic decrease of film thickness is observed with decreasing ink load as the luminosity, L, varies from 0 to 50. Films were discontinuous below L ¼ 70. Sheet resistivities of inkjet-printed films and spin-coated films were measured using the four-point probe method and results are shown in Figure 3.8. Two conclusions can be drawn from these results. First, as expected, a distinct increase in sheet resistivity is observed with decreasing film thickness. A change in sheet resistivity of one order of magnitude is observed as the thickness changes from 100 to 200 nm. It should be noted that changes in
Sheet resistivity (ohms/sq.)
10,000
L = 50 L = 30 L = 20
Spincoated
L = 10
Inkjetted
L=0
1000 0
50
100
150
200
250
Thickness (nm)
FIGURE 3.8 Measured PEDOT–PSS sheet resistivity of inkjet-printed and spin-coated films with various thicknesses. The measurement was carried out in air, at room temperature, using a four-point probe.
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3-7
sheet resistivity with thickness are more drastic for films thinner than 200 nm. Second, the sheet resistivities of spin-coated films were slightly higher than those of inkjet-printed films. We suspect that this is due to the difference in drying processes between the two fabrication techniques. In spincoating, the film dries rapidly during the rotation of the substrate and films may be formed with a random coil conformation with a large degree of disorder. This initial fast drying is even more apparent in thinner films. On the other hand, the drying of inkjet-printed films is much slower than those of spincoated counterparts; thus there is greater time for equilibration of the material during the drying process. This possibly creates a more favorable (homogeneous) distribution of PEDOT–PSS polymer chains, resulting in the lower sheet resistivity. Surface morphology of films was measured by the optical interferogram (Wyko) of 60 mm 45.5 mm area. For comparison, formulated ink was spin-coated on a glass substrate with the same thickness of 220 nm, which is similar to that of an inkjet-printed layer. No pin-holes were observed on layers deposited using either techniques. However, in both cases, many randomly positioned voids have been observed on the dried PEDOT–PSS surface. Figure 3.9a shows the surface of the inkjet-printed film
(a) Inkjet printed and dried at 110⬚C,
nm
27.8
20.0 15.0 10.0 5.0 0.0
59.7
−5.0 −10.0
45.5 µm
(b) Spin-coated and dried at 110⬚C,
−15.0
nm
39.1 30.0 25.0 20.0 15.0 10.0 5.0
59.7
0.0 −5.0 −10.0 −15.0 45.5 µm
−20.0 −25.0 −30.0
FIGURE 3.9 Surface topography of printed films of the PEDOT–PSS. (a) Inkjet printed and dried at 1108C and (b) spin-coated and dried at 1108C. Each layer was dried on the hot plate in atmospheric condition for 2 h.
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when the substrate was heated at 1108C. The pit diameter was 2–3 mm and the depth was approximately 25–35 nm (Ra ¼ 3.89 nm). The spin-coated PEDOT–PSS film showed smaller pits than that of the inkjet-printed film. The depth of voids was about 15 nm with a width of 1–2 mm and roughness average of 4.29 nm (Figure 3.9b). Since the miscibility of PEDOT–PSS and glycerol is poor, as water evaporates, the remaining glycerol coalesces onto the PEDOT–PSS surface and leaves behind voids in the PEDOT–PSS layer upon evaporation. Smaller pits were observed as the layer thickness decreases since there is less glycerol on the PEDOT–PSS surface. Compared to inkjet printing, the centrifugal effect in spin-coating shortens the drying time of the two solvents, thus resulting in smaller pits than the case of inkjet. In order to investigate the referenced inkjet-printed film in an OLED, some inkjetted PEDOT–PSS films were used as the anode. On top of the inkjet-printed anode, the hole transport layer (HTL) solution (TPD, [N 0 ,N 0 ,-bis(3-methylphenyl)-N 0 ,N 0 dimethyl benzidine] 67.6 wt.%, polycarbonate (PC) 29.0 wt.%, rubrene 3.4 wt.%, 10.35 mg=ml chloroform) was spin-coated at 1000 rpm for 1 min in a class 100 cleanroom. A 60 nm layer of tris-(8-hydroxyquinoline)-aluminum (Alq3) was then thermally deposited under the high vacuum at the rate of 0.7 A˚=s. Then, a 300 nm layer of Mg:Ag (magnesium:silver) was thermally coevaporated at the ratio of 10:1 on the top of electron transport layer (ETL) layer (Figure 3.10). The thickness of spin-coated layers was matched to one of the inkjet-printed layers (L ¼ 0). Additionally, Figure 3.10 shows results of OLED characteristics with the same layer configuration except that ITO is used as the anode layer. Figure 3.11 shows the current density (mA=cm2) vs. bias voltage (V) for the various devices made with varying L values. As expected, a decrease in current density was observed with decreasing thickness (increasing L) and increasing sheet resistivity of the PEDOT–PSS layer. The forward light output (Figure 3.12) from the devices followed the same trends. Lower light output at a given voltage can be seen with decreasing thickness (increasing sheet resistivity). At any given voltage, the device with the spin-coated layer had a higher current and brightness compared to an OLED fabricated with inkjet-printed anode layer and this is built due to the slight difference in sheet resistivity of spin-coated and inkjet-printed layers. The peak external quantum efficiencies (hext) (%) of devices containing PEDOT–PSS anodes were compared with those of devices utilizing an ITO anode, at approximately 1.2–1.3% efficiency (Figure 3.13). These values were higher than that of devices made by spin-coating, which had 0.9% efficiency. These results show that the surface morphology of the anode affects the hext of OLEDs. Values of hext of inkjet-printed PEDOT–PSS anode remained constant or demonstrated only slight changes at higher voltages, unlike the hext of ITO anode, which rapidly dropped off after the peak hext is achieved. This was presumably due to the low current density of PEDOT–PSS anodes and the topography of their surface. However, more detailed studies will be necessary to explore the precise relation between PEDOT–PSS surface morphology and OLED performances.
Mg:Ag 300 nm
Alq3:60 nm
TPD:rubrene:PC
PEDOT–PSS
FIGURE 3.10
The device configuration of OLED.
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Inkjet Printing and Patterning of PEDOT–PSS: Application to Optoelectronic Devices
25
Inkjetted, t = 220 nm Spin-coated, t = 220 nm ITO
20
Current density (mA/cm2)
Current density (mA/cm2)
25
15
10 5
L = 10, t = 200 nm L = 20, t = 191 nm L = 30, t = 160 nm L = 50, t = 125 nm ITO
20 15 10 5
0
0 0
5
10
15
20
0
5
Voltage (V)
FIGURE 3.11
1000
Inkjetted, t = 220 nm Spin-coated, t = 220 nm ITO
800 600 400 200
600 400 200 0
0
FIGURE 3.12
5
10 Voltage (V)
15
0
20
5
10
15
20
Voltage (V)
Forward light output vs. bias voltage of devices having various PEDOT–PSS thicknesses.
1.6 Inkjetted Spin-coated ITO
0
5
External quantum efficiency (%)
External quantum efficiency (%)
20
L = 10, t = 200 nm L = 20, t = 191 nm L = 30, t = 160 nm L = 50, t = 125 nm ITO
800
0
2.4 2.2 2.0 1.8 1.6 1.4 1.2 1.0 0.8 0.6 0.4 0.2 0.0 −0.2
15
Current density vs. bias voltage of devices having various PEDOT–PSS thicknesses.
Forward light output (cd/m2)
Forward light output (cd/m2)
1000
10
Voltage (V)
10 Voltage (V)
FIGURE 3.13
15
20
1.4
L = 10 L = 20 L = 30 L = 50 ITO
1.2 1.0 0.8 0.6 0.4 0.2 0.0
−0.2
0
5
10
15
20
Voltage (V)
External quantum efficiencies vs. bias voltage of devices having various PEDOT–PSS thicknesses.
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FIGURE 3.14
Photographs of OLEDs with inkjet deposited PEDOT–PSS anode on glass substrates.
Figure 3.14 shows a simple OLED logo designed on graphical software and was printed onto the glass substrate of 2.54 2.54 cm2 using inkjet printing and PEDOT–PSS-based ink. The thickness of the anode can be controlled simply by using HSL color function.
3.3
Deactivatived Patterning of PEDOT–PSS
In this section, we demonstrate the approach to modify the sheet resistivity of PEDOT–PSS anode by controlling the desktop inkjet printer using HSL color functions and oxidizing agents, in order to fabricate gray-scale electroluminescent images via single step process. We would like to stress the fact that we are patterning the anode in a continuous fashion. Our aim here is to obtain stationary displays (signs and logos) rather than dynamic ones (active or passive matrix) where the user can control the picture elements and colors. Two types of oxidants were used. One of the inks consisted of an oxidizing agent in the form of 2 wt.% of sodium hypochlorite aqueous solution, and a surfactant of 0.13 wt.% and the other ink contained approximately 50 wt.% of hydrogen peroxide. The ink was loaded into the black inkjet cartridge of the printer. The viscosity and surface tension of ink were optimized to obtain the best image intensity and resolution. The oxidizing agent interacts with the polymeric anode and increases its sheet resistivity. The extent of this chemical reaction was controlled through the adjustment of the ink volume dispensed and the postprinting processing treatment. During the postprinting processing, the substrates were moderately (408C–608C) heated to enhance the chemical reaction between the PEDOT–PSS and the oxidizing agent. In the case of sodium hypochlorite ink, to end this reaction, the printed ink needs to be thoroughly washed away with DI water, whereas hydrogen peroxide ink does not require the washing process. The substrates were then
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TABLE 3.1 HSL Values and Colors Used in Programing the Ink Loads (S ¼ 0 and H was Not Defined) Color
No.
Luminosity (L)
#1
255
#2
230
#3
215
#4
200
#5
170
#6
140
#7
125
#8
110
#9
50
# 10
0
dried and the organic layers were deposited thereafter. The ink load of the printing can be controlled through the HSL function as mentioned above. Controlling the dot density of the oxidizing agent (along with careful postprinting processing) allows accurate controlled reaction with the polymeric anode. It can be controlled by the ink load and spatial density (at a fixed solution concentration) of the inkjet droplets (darkness) over the substrate area. Luminosity (L), which controls the amount of printed oxidative ink, was changed over 10 different values as shown in Table 3.1. The extent of the chemical reaction between oxidizing agent and PEDOT–PSS can be controlled through the adjustment of the volume of the dispersed oxidative ink, and resultant sheet resistivities of PEDOT–PSS with corresponding luminosity values were measured by a four-point probe method. Figure 3.15 shows the sheet resistivity as a function of the processing time of sodium hypochlorite ink at a given value of L ¼ 0. As apparent from the figure, the most significant change in PEDOT–PSS sheet
Sheet resistivity (ohms/sq.)
1.E+06
1.E+05
1.E+04
1.E+03
1.E+02 0
50
100
150
200
Time (min)
FIGURE 3.15 Measured PEDOT–PSS sheet resistivity of inkjet-printed sheets with time (at L ¼ 0, sodium hypochlorite ink). The measurement was carried in air, at room temperature, and using a four-point probe.
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resistivity occurs within the first few minutes of the oxidizing agent being deposited on the surface. Further oxidation of PEDOT–PSS is achieved at longer time periods. A possible explanation is that initially a rapid reaction occurs at the interface of the oxidizing agent and PEDOT–PSS layer and then the remaining oxidizing agent slowly diffuses and reacts with the inner regions of the anode layer, causing further sheet resistivity change. This diffusion process can be enhanced by moderately heating the substrate. Moreover, at a fixed processing time after printing the oxidizer, a systematic increase of sheet resistivity is observed with increased ink load (Figure 3.16). A change in sheet resistivity of three orders of magnitude is observed as L varies from 255 (white, no ink) to 0 (black, maximum allowed ink). On the other hand, the reaction between hydrogen peroxide and PEDOT–PSS shows a slightly different behavior. Unlike sodium hypochlorite, hydrogen peroxide has a lower boiling point (b.p. ¼ 1268C in 50 wt.% aqueous solution) and at room temperature, most of the ink evaporates within 20 min. One of the advantages of using hydrogen peroxide ink is that it does not require the washing process, thus reducing the opportunity for the unprinted polymer area to be exposed to water. In addition, hydrogen peroxide is less reactive to the metal parts of the ink cartridges. It can reduce the risk of damaging the printheads and ease the printing process. Figure 3.17 shows measured sheet resistivity of the resulted films that were dried at three different temperatures (208C, 408C, and 508C) with different L values. Moderately heating the substrate increases the speed of solvent evaporation and additionally enhances the chemical reaction. This was because at lower temperatures, the oxidant stayed on the substrates longer and allowing the reaction to proceed at a steady rate. In all cases, systematic increase of sheet resistivity is observed with lowering the L values. For the OLED fabrication, the HTL was deposited on top of the PEDOT–PSS anode by spincasting of a layer of TPD and PC, 1 wt.% concentration inchloroform, to a thickness of 55–60 nm. The relative percentage of TPD and PC was 70% and 30%, respectively. For the ETL, Alq3 was thermally sublimed at a rate of 1.0 A˚=s to a thickness of 60 nm under high vacuum (106 to 107 torr). Both the HTL and the ETL covered the entire area of the substrates. Mg and Ag were coevaporated over the entire ETL in a ratio of 10 (Mg) to 1 (Ag) to form a 300 nm thick cathode. The device structure of these OLEDs is shown in Figure 3.18. The devices were tested in a nitrogen glove box.
Sheet resistivity (ohm/sq.)
1.E+06
1.E+05
1.E+04
1.E+03
1.E+02 250
200
150
100
50
0
Luminosity
FIGURE 3.16 Measured sheet resistivity for various luminosity values: the reaction between PEDOT–PSS and sodium hypochlorite ink. The measurement was carried in air, at room temperature, and using a four-point probe.
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1.E+07
Sheet resistivities (ohms/sq.)
1.E+06 1.E+05 1.E+04 1.E+03 1.E+02 1.E+01 Dried at 50°C Dried at 40°C Dried at 20°C 1.E+00 250
200
150
100
50
0
Luminosity
FIGURE 3.17 Measured sheet resistivity for various luminosity values and temperatures: the reaction between PEDOT–PSS and hydrogen peroxide ink. The measurement was carried in air, at room temperature, and using a four-point probe.
Only the results of OLEDs fabricated using anodes, sodium hypochlorite-treated PEDOT:PSS are presented here. Upon applying the bias voltage across the device, we anticipate a part of this voltage drop across the resistive anode path. The higher the sheet resistivity, the lower the current. Figure 3.19 shows the current density (mA=cm2) vs. bias voltage (V) for the various devices made with various L values. A decrease in current density is observed with increasing sheet resistivity (decreasing L) of the PEDOT– PSS layer. Since the light output is proportional to the current, the forward light output from the devices also follows a trend similar to the current density (Figure 3.20). For a given voltage, lower light output is observed as the sheet resistivity is increased. For example, at 10 V, the light output levels were measured to be 160, 107, 61, 36, and 24 cd=m2 for devices printed with the luminosity parameter L preset to 255 (0%), 230 (10%), 200 (22%), 140 (45%), and 110 (57%), respectively. This variation in brightness due to different sheet resistivity values for a given voltage introduces the potential for gray-scale imaging, thus allowing us to print any graphical or photographic image into an active light-emitting device. The peak external quantum efficiency (%) remains relatively constant at about 0.5% (Figure 3.21), but shifts to higher values with increasing sheet resistivity. Although our materials are not purified and the
Patterned
PEDOT/PSS ANODE HTL ETL CATHODE
TPD:PC Alq3 (60 nm) Mg:Ag (300 nm)
FIGURE 3.18 The schematic image of the fabricated OLEDs where the PEDOT–PSS was patterned by inkjet printing. There was no patterning of the HTL, ETL, or cathode.
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60 L = 255 L = 230
50
L = 215
Current density (mA/cm2)
L = 200 L = 140
40
L = 110 L = 50
30
20
10
0 0
5
10
15
20
25
Voltage (V)
FIGURE 3.19 Current density vs. bias voltage of devices having various PEDOT–PSS sheet resistivities (treated by sodium hypochlorite).
450 L = 255
Forward light output (cd/m2)
400
L = 230
350
L = 21
300
L = 140
L = 200 L = 110
250
L = 50
200 150 100 50 0 −50 0
5
10
15
20
25
Voltage (V)
FIGURE 3.20 Forward light output vs. bias voltage for OLEDs made with PEDOT–PSS having various sheet resistivities (treated by sodium hypochlorite).
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Inkjet Printing and Patterning of PEDOT–PSS: Application to Optoelectronic Devices
0.6
L = 255
External quantum efficiencies (%)
L = 230 L = 215
0.5
L = 200 L = 140 L = 110
0.4
L = 50 0.3
0.2
0.1
0.0
0
5
10
15
20
25
Voltage (V)
FIGURE 3.21 External quantum efficiencies vs. bias voltage for OLEDs made with PEDOT–PSS having various sheet resistivities (treated by sodium hypochlorite).
processing is not optimized, the peak value above compares well with the 0.7% efficiency values reported by Kim et al. [3]. One of the proposed chemical mechanisms of sodium hypochlorite treatment is shown in Figure 3.22. In this reaction, sodium hypochlorite is the oxidant and converts thiophene (A) in PEDOT [43] to its corresponding thiophene-1-oxide (B). Spontaneously, thiophene-1-oxide (B) is converted to its corresponding thiophene-1,1-dioxide (C). The thiophene-1-oxide (B) is presumed to be a reaction intermediate to thiophene-1,1-dioxide (C) [43,44]. Finally, further oxidation of thiopene1,1-dioxide (C) causes the extrusion of SO2 from this compound and the attachment of hydroxyl groups due to the nucleophilic attack by water resulting in structure (D). In Figure 3.23, the FT-IR spectra of the sodium hypochlorite PEDOT–PSS film and the pristine PEDOT–PSS film are compared over a range of wavenumbers from 500 to 4000 cm1. Minor differences in relative band intensities are observed. In the sodium hypochlorite–treated film, the absorbance at 1343 and 1132 cm1 results from the asymmetric and symmetric stretching vibrations of the SO2 group, respectively [44–46]. The absorbance at 1044 cm1 is due to the symmetrical stretching vibration of the sodium sulfonate (SO3) group of PSS-Na. To observe the completion of this oxidation reaction,
(A)
O
(B)
O
* S
O
NaOCl * n
H2O
O
* S O
FIGURE 3.22
(C)
O
NaOCl * n
H2O
(D)
O
* O
S
O
O
NaOCl * n
H2O
Chemical mechanism of the PEDOT oxidation with sodium hypochlorite.
O
*
* OH OH Na2SO4
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Absorbance
Conjugated Polymers: Processing and Applications
(a)
Sodium hypochlorite–treated PEDOT–PSS
(b)
Pristine PEDOT–PSS
(a)
(b)
4000
3500
3000
2500
2000
Wave number (cm
1500
1000
500
−1)
FIGURE 3.23 FT-IR spectra of (a) inkjet–treated (oxidized by sodium hypochlorite) PEDOT–PSS film and (b) pristine PEDOT–PSS film on the plastic substrates.
Absorbance
structure D was synthesized by mixing the excess molar amount of NaOCl aqueous solution with the diluted PEDOT–PSS dispersion (Baytron P, Bayer). The solution was allowed to dry and the resultant powders were analyzed by the FT-IR along with pristine PEDOT–PSS powders. Figure 3.24 shows the FT-IR spectra of both powders. Distinct differences are observed in relative band intensities. The bands
(a)
Sodium hypochlorite–treated PEDOT–PSS
(b)
Pristine PEDOT–PSS
(a)
(b)
4000
3500
3000
2500
2000
Wave number
1500
1000
500
(cm−1)
FIGURE 3.24 FT-IR spectra of (a) fully oxidized (by sodium hypochlorite) PEDOT–PSS structure and (b) pristine PEDOT–PSS structure with KBr pellets. Dried materials (powders) were pressed into the KBr pellets and analyzed. In both cases, 128 scans were taken with a resolution of 4 cm1.
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(A)
(B)
O
O
* S
* n
O
O
H2O2
O
O
O
O
&
H2O *
S O
FIGURE 3.25
3-17
O
* n
*
S O
O
* n
Chemical mechanism of the PEDOT oxidation with hydrogen peroxide.
Absorbance
due to the hydroxyl group (O–H) stretching and bending vibrations, and C–O stretching vibrations are found at 3493, 1410, and 1201 cm1, respectively. The absorbance at 1616 cm1 is due to the stretching vibrations (C==C) of the conjugated olefin (unsaturated open-chain hydrocarbons). The band at 1134 cm1 results from the stretching vibrations of SO4 in the sulfate ion (SO42) [46]. Finally an additional band at 974 cm1 is seen due to the vibrations of the O–Cl group in the remaining sodium hypochlorite. Furthermore, the possible chemical pathway of the hydrogen peroxide oxidation of PEDOT–PSS is proposed in Figure 3.25. The suggested reaction route is similar to that of sodium hypochlorite, which leads the thiophene compound to the corresponding sulfoxide (SO2) [47]. In addition, it is possible to degrade oxyethylene rings upon the exposure to oxidizing species. The suggested structures include the formation of glycol diformate (–O–C==O) along with the sulfoxide [48]. The comparison of FT-IR spectra of hydrogen peroxide–treated PEDOT–PSS and pristine PEDOT– PSS is shown in Figure 3.26. To ensure the complete oxidation, the PEDOT–PSS (Baytron P) films were originally prepared on the silicon substrates and excess amount of hydrogen peroxide aqueous solution was applied to synthesize structure (B). In relative band intensities, the distinct differences for the pristine PEDOT–PSS at 1309, 1084, and 931 cm1 were probably due to the suflonate (SO3) group of
(a)
Hydrogen peroxide–treated PEDOT–PSS
(b)
Pristine PEDOT–PSS
(a)
(b)
4000
3500
3000
2500
2000
1500
1000
500
Wave number (cm−1)
FIGURE 3.26 FT-IR spectra of (a) fully oxidized (by hydrogen peroxide) PEDOT–PSS structure and (b) pristine PEDOT–PSS structure.
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0%
(a)
33%
(b)
µm
µm 2 1
2 1
RMS = 9.281 nm
(c)
RMS = 6.679 nm (d)
80%
100%
µm
µm
2
2
1
1
RMS = 5.337 nm
RMS = 3.193 nm
FIGURE 3.27 Atomic force microscopy (AFM) images of PEDOT–PSS layers with different sets of L values. (a) L ¼ 255, (b) L ¼ 170, (c) L ¼ 50, and (d) L ¼ 0 (3 3 mm2).
100 90
Transmittance,T,%
80 70
(a)
(b) (c)
60 50 40
(a)
Orgacon substrate
30
(b)
L = 0 with hydrogen peroxide
20
(c)
L = 0 with sodium hypochlorite
10 0 400
450
500
550
600 650 Wavelength (nm)
700
750
800
FIGURE 3.28 Optical absorption spectra of (a) pristine PEDOT–PSS, (b) hydrogen peroxide–treated PEDOT–PSS, and (c) sodium hypochlorite–treated PEDOT–PSS.
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3-19
PSS, dissociated from PEDOT after the oxidation reaction. In the hydrogen peroxide-treated film, the absorbance at 1360 and 1182 cm1 are the results from the vibration of the SO2 group. The bands appeared at 1751, 1678, and 1416 cm1 corresponding to the formation of the glycol diformate (–O–C==O) structure. Atomic force microscopy (AFM) was used to image the surface and highlight the morphological changes between the pristine PEDOT–PSS surface (L ¼ 255) and the sodium hypochlorite–treated surfaces (L ¼ 170, 50, and 0) (Figure 3.27). The root-mean-square (RMS) roughness of the surface decreased from 9.3 to 3.2 nm as the values of L decreased from 255 to 0. Although changes in PEDOT–PSS morphology were observed, no distinct film thickness difference between conductive and nonconducting (oxidized) surfaces was detected by the step profilometer. The optical transmission spectra of sodium hypochlorite–treated, hydrogen peroxide–treated, and pristine PEDOT–PSS show only small differences in light absorption in the visible area (Figure 3.28). Unlike the electron transfer reaction in electrochromic materials [44,50], the oxidization reaction here only slightly darkens the PEDOT–PSS layer. We suspect that in our case, the oxidation process does not occur all throughout the film but is rather confined to a limited depth that cannot be determined accurately at this time [51]. The technique discussed above can be easily extended to print any pattern including photographs. Figure 3.29 shows a printed electroluminescent image of the sun on a flexible plastic substrate. It is worth emphasizing that all layers were deposited over the entire surface of the substrate, without any patterning.
FIGURE 3.29 Photographs of gray-scale OLEDs that were patterned on PEDOT–PSS surface by inkjet printer on plastic substrates. Conductivities were modified to exhibit the contrasts.
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Conclusion
The strategy of controlling ink loads using HSL color function and a black cartridge has been described. This approach can be used to create a library of electrodes with different sheet resistivities within a few minutes. In addition, PEDOT–PSS morphology and thickness can be greatly influenced by the processing conditions, additives, and solute concentrations. Surface roughness and conductivity of inkjetted films can be comparable to spin-coated ones. This technique can be easily extended to print conductive polymers onto more delicate surfaces, such as textiles or paper-based substrates. The use of inkjet techniques to control the sheet resistivity of conducting polymers can be beneficial in the gray-scale design of OLEDs, as well as in numerous other applications where the desired values of resistivities cannot be achieved using other techniques in faster time and cost-effective manner.
Acknowledgment We would like to thank Dr. Dino Pardo for his comments and fruitful talks.
References 1. Aleshin, A.N., S.R. Williams, and A.J. Heeger. 1998. Synth Met 94:173. 2. Kirchmeyer, S., and K. Renter. 2005. J Mater Chem 15:2077. 3. Kim, W.H., A.J. Ma¨kinen, N. Nikolov, R. Shashindhar, H. Kim, and Z.H. Kafafi. 2002. Appl Phys Lett 80:3844. 4. Kim, W.H., G.P. Kushto, H. Kim, and Z.H. Kafafi. 2003. J Polym Sci 41:2522. 5. Ghosh, S., and O. Ingana¨s. 2001. Synth Met 121:1321. 6. Lai, S.L., M.Y. Chen, M.K. Fung, C.S. Lee, and S.T. Lee. 2003. Mater Sci Eng B 104:26. 7. Timpanaro, S., M. Kemerink, F.J. Touwslager, M.M. De Kok, and S. Schrader. 2004. Chem Phys Lett 394:339. 8. Jo¨nsson, S.K.M., J. Birgerson, X. Crispin, G. Greczynski, W. Osikowicz, A.W. Denier van der Gon, W.R. Salaneck, and M. Fahlman. 2003. Synth Met 10361:1. 9. Louwet, F., L. Groenendaal, J. Dhaen, J. Manca, J. Van Luppen, E. Verdonck, and L. Leenders. 2003. Synth Met 135:115. 10. Ouyang, J., C.-W. Chu, F.-C. Chen, Q. Xu, and Y. Yang. 2005. Adv Funct Mater 15:203. 11. Huang, J., P.F. Milelr, J.C. de Mello, A.J. de Mello, and D.D.C. Bradley. 2003. Synth Met 139:569. 12. Jukes, P.C., S.J. Martin, A.M. Higgins, M. Geoghegan, R.A.K. Jones, S. Langridge, A. Wehrum, and S. Kirchmeyer. 2004. Adv Mater 16:807. 13. Huang, J., P.F. Miller, J.S. Wilson, A.J. de Mello, J.C. de Mello, and D.D.C. Bradley. 2005. Adv Funct Mater 15:290. 14. Kim, J.Y., J.H. Jung, D.E. Lee, and J. Joo. 2002. Synth Met 126:311. 15. Lee, C.S., J.Y. Kim, D.E. Lee, Y.K. Koo, J. Joo, S. Han, Y.W. Beag, and S.K. Koh. 2003. Synth Met 135–136:13. 16. Aernouts, T., P. Vanlaeke, W. Geens, J. Poortmans, P. Heremans, S. Borghs, R. Mertens, R. Andriessen, and L. Leenders. 2004. Thin Solid Films 451–452:22. 17. Hohnholz, D., H. Okuzaki, and A.G. MacDiarmid. 2005. Adv Funct Mater 15:51. 18. Zhang, F., T. Nyberg, and O. Ingana¨s. Nano Lett 2:1373. 19. Tan, L., Y.P. Kong, S.W. Pang, and A.F. Yee. 2004. J Vac Sci Technol B 22:2486. 20. Nilsson, D., M. Chen, T. Kugler, T. Remonen, M. Armgarth, and M. Berggren. 2002. Adv Mater 14:51. 21. Yoshioka, Y., P.D. Calvert, and G.E. Jabbour. 2005. Macromol Rapid Commun 26:238. 22. Calvert, P. 2001. Chem Mater 13:3299. 23. Calvert, P., Y. Yoshioka, and G.E. Jabbour. 2004. Learning from nature how to design new implantable biomaterials: From biomineralization fundamentals to biomimetic materials and processing routes. Dordrecht, Netherlands: Kluwer Academic, p. 169.
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24. 25. 26. 27. 28. 29. 30. 31. 32. 33. 34. 35. 36. 37. 38. 39. 40. 41. 42. 43. 44. 45. 46. 47. 48. 49. 50. 51.
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de Gans, B.-J., P.C. Duineveld, and U.S. Schubert. 2004. Adv Mater 16:203. Bharathan, J., and Y. Yang. 1998. Appl Phys Lett 72:2660. Chang, S.-Ch., J. Liu, J. Baharathan, Y. Yang, J. Onohara, and J. Kido. 1999. Adv Mater 11:734. Kobayashi, H., S. Kanabe, S. Seki, H. Kigchi, M. Kimura, I. Yudasaka, S. Miyashita, T. Shimoda, C.R. Towns, J.H. Burroughes, and R.H. Friend. 2000. Synth Met 111:125. Shibusawa, M., M. Kobayashi, J. Hanari, K. Sunohara, and N. Ibaraki. 2003. IEICE Trans Electr E86:2269. Shimoda, T., K. Morii, S. Seki, and H. Kiguchi. 2003. MRS Bull 28:821. Sirringhaus, H., T. Kawase, R.H. Friend, T. Shimoda, M. Inbasekaran, W. Wu, and E.P. Woo. 2000. Science 290:2123. Liu, Y., T. Cui, and K. Varahramyan. 2003. Solid State Electron 47:1543. Liu, Y. and T. Cui. 2005. Macromol Rapid Commun 26:289. Mo¨ller, M., S. Asaftei, D. Corr, M. Ryan, and L. Walder. 2004. Adv Mater 16:1558. Mabrook, M.F., C. Pearson, and M.C. Petty. 2005. Appl Phys Lett 86:013507. Setti, L., A. Fraleoni-Morgera, B. Ballarin, A. Filippini, D. Frascaro, and C. Piana. 2005. Biosens Bioelectr 20:2019. Ballarin, B., A. Fraleoni-Morgera, D. Frascaro, S. Marazzita, C. Piana, and L. Setti. 2004. Synth Met 146:201. Deegan, R.D., O. Bakajin, T.F. Dupont, G. Huber, S.R. Negel, and T.A. Witten. 1997. Nature 389:827. Fischer, B.J. 2002. Langmuir 18:60. Deegan, R.D. 2000. Phys Rev E 61:475. Deegan, R.D., O. Bakajin, T.F. Dupont, G. Huber, S.R. Nagel, and T.A. Witten. 2000. Phys Rev E 62:756. Steiger, J., S. Heun, and N. Tallant. 2003. J Imag Sci Tech 47:473. Maenosono, S., C.D. Dushkin, S. Saita, and Y. Yamaguchi. 1999. Langmuir 15:957. Lin, S., and R.M. Carlson. 1984. Environ Sci Technol 18:743. Jiang, B., and T.D. Tilley. 1999. J Am Chem Soc 121:9744. Weiss, R.A., A. Sen, C.L. Willis, and L.A. Pottick. 1999. Polymer 32:1867. Socrates, G. 1980. Infrared characteristics group frequencies. New York: John Wiley & Sons. Hulea, V., F. Fajula, and J. Bousquet. 2001. J catal 198:179. Maurino, V., P. Calza, C. Minero, E. Pelizzetti, and M. Vincenti. 1997. Chemosphere 35:2675. Heuer, H.W., R. Wehrmann, and S. Kirchmeyer. 2002. Adv Func Mater 12:89. Argun, A.A., A. Cirpan, and J.R. Reynolds. 2003. Adv Mater 15:1338. Zotti, G., S. Zecchin, G. Schiavon, F. Louwet, L. Groenendaal, X. Crispin, W. Osikowicz, W. Salaneck, and M. Fahlman. 2003. Macromolecules 36:3337.
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4 Printing Organic Electronics on Flexible Substrates 4.1
Introduction......................................................................... 4-1 Applications for Printed Organic Electronics
4.2 4.3
Nathaniel D. Robinson and Magnus Berggren
4.1
.
Coating with
Patterning Electroactive Materials ................................... 4-10 Additive Patterning Techniques
4.5 4.6 4.7
Market
Substrates for Printed Organic Electronics ....................... 4-5 Coating Processes ................................................................ 4-8 Spin-Coating . Linear Coating Processes Flow Processes
4.4
.
.
Subtractive Patterning
.
Other Novel
Curing................................................................................. 4-23 Encapsulation .................................................................... 4-24 Conclusion ......................................................................... 4-24
Introduction
Aside from versatility, one of the most cited reasons for using organic electronic materials is that they are solution processable, offering reduced costs and the possibility for roll-to-roll and other high-volumeprocessing techniques. Over the past 15 years, solutions and suspensions of conjugated polymers have been developed to the point where they are manufacturable, as demonstrated by products such as the organic light-emitting diodes (OLEDs) used in Philips’ electric razor. The intention of this chapter is to illustrate many of the printing techniques available for the production of organic electronics, with a focus on paper and flexible substrates, and highlight some of the challenges. The motivation for using organic electronic materials is twofold. First, these materials offer functionality not available from traditional materials, such as the colors available from OLEDs [1]. Second, organic materials (polymeric or molecular) often lend themselves to solution processing— processing in a fluid (or paste) form—with the same techniques used to print newspapers, magazines, or even desktop-published documents. Printing techniques, especially when used in a roll-to-roll manufacturing process, offer tremendous potential cost savings over clean room processing of crystalline materials. To manufacture electronic components, one generally must be able to produce conducting lines, insulating lines (and layers, if multilayer devices are to be constructed), and electronically active lines (channels in transistors, etc.). Auxiliary features, such as encapsulation to protect the device from its surroundings or electrolytes in electrochemical devices, may also be required. Each layer typically 4-1
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involves one or more steps, although clever processing has been demonstrated to combine the formation of different structures simultaneously.
4.1.1
Applications for Printed Organic Electronics
The first products based on conjugated polymers have been OLEDs and displays based on arrays of such devices, possibly because they were some of the first devices created in research laboratories, and they offer colors [1] and flexibility that are not available from traditional inorganic materials. Thus, improvements to materials and structures (typically layers of different materials) are coming at an alarming rate. Products containing OLEDs already on the market include an electric razor from Philips and a digital camera from Kodak. A prototype OLED display from DuPont OLED Displays appears in Figure 4.1. Another potentially large application for organic electronics is the field-effect transistor and circuits thereof. However, these devices have not widely reached commercial production because they offer little advantage over their silicon counterparts. It is likely that transistors will first be widely produced in conjunction with a display technology, such as in active-matrix light-emitting diode (LED) screens. However, efforts to produce inexpensive printed radio frequency identification (RFID) tags have produced impressive results. For example, the German company PolyIC [2] has published several results including a 4-bit transponder chip [3]. Reflective displays, such as the electrochromic poly(3,4-ethylenedioxythiophene)–poly(styrene sulfonate) (PEDOT:PSS) [4] matrix-addressed display shown in Figure 4.3b [5], are another application for organic electronic materials. These displays reflect or absorb light from the environment, rather than emit light, requiring much less power under operation. There are three basic technologies for making reflective diplays: electrochromism [6,7], electrophoresis [8], and electrowetting [9]. Electrochromism describes a color change as a result of an applied potential, typically via an electrochemical mechanism. Many electrochemically switchable materials such as polythiophenes and polyaniline have this capability. Table 4.1 shows a very few of the materials and colors available today. Other nonpolymeric materials such as viologens [10] offer a higher contrast, but are not themselves conducting and therefore must be combined with a conducting polymer or another conducting
FIGURE 4.1 (See color insert following page 8-22.) 14.1 in. diagonal, solution processed, amorphous silicon active-matrix OLED display fabricated by DuPont OLED Displays with active matrix backplane provided by SEC. Full-color displays based on OLEDs have brilliant colors, have a viewing angle approaching 1808, and are much more efficient than the LCD displays currently used in most applications.
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Printing Organic Electronics on Flexible Substrates TABLE 4.1 Examples of Electrochromic Conjugated Polymers Name
Chemical Structure
Poly(3,4-ethylenedioxythiophene) (PEDOT)
Colors Oxidized: faint blue
S
Reduced: deep blue
O
O
H
Poly(aniline) (PANI)
Neutral: colorless
N
Notes Highly conductive when oxidized (even partially) Doped state is especially stable Commercially available in many forms, including aqueous suspensions Commercially available in many forms
Oxidized 1: green Not conducting in Protonated oxidized 1: blue neutral state Oxidized 2: blue Protonated oxidized 2: violet
Poly(3-hexylthiophene) (P3HT)
S
Oxidized: clear Neutral: red-pink
Commercially available Not conducting in neutral state Other alkyls can be substituted for hexyl group Need at least butyl for polymer to be soluble
material to be used in a display. Of course, these nonpolymeric materials are useful for more than their colors. Many are used in printed electronic components such as field-effect transistors in RFID tags, for example. Electrochromic displays are simple electrochemical cells with an electrochromic material as one of the electrodes. A potential applied between the anode and cathode, typically less than 5 V, produces the desired color change. Once switched, the current between the anode and cathode diminishes. Most electrochromic displays have a memory (stay switched even if no potential is applied as long as the electronic connection between anode and cathode is broken), many on the order of hours. Thus, they are ideal for low-power applications. Unfortunately, many printed electrochromic devices are relatively slow. Speeding up the device requires a liquid electrolyte, which is difficult to manufacture and encapsulate in an inexpensive device. All-printed electrochromic displays combined with simple electrochemical transistor logic are presently developed at Linko¨ping University and Acreo AB [11]. The goal is a true electronic paper technology. PEDOT:PSS coated on photographic quality paper is subtractively ‘‘printed’’ to pattern nonconducting lines within the polymer film. Electrolyte is then screen-printed on top of the patterned PEDOT:PSS to form lateral electrochemical devices as presented in Figure 4.2. From this, active-matrixaddressed displays and simple logic circuits have been realized. Electrophoretic displays, such as those produced by E-Ink [8], are based on two immiscible fluids, or a fluid and solid particles, which are displaced by the electric field within an encapsulated cell. An applied
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Visible electrode
Opaque electrolyte
Counter electrode (a)
Transparent electrolyte
Counterelectrode
Visible electrode
(b)
FIGURE 4.2 Reflective electrochromic display structures. (a) Vertical display structure cross section, to be seen from the top. The opaque (white) electrolyte hides the counterelectrode. (b) Lateral display structure cross section. The transparent electrolyte allows the visible electrode to be seen. The counterelectrode should be placed under another printed object if it must be hidden.
field drives the positively charged fluid (or particles) towards the negatively charged electrode and vice versa. The display is observed through one of the electrodes (which must be transparent). Unlike electrochromic displays, electrophoretic displays require that a potential be applied as long as the display is active, but the power requirement is still much smaller than for luminescent displays. Since electrophoretic displays require a fluid bulk and a transparent electrode, they typically are ‘‘printed’’ on an indium tin oxide (ITO) substrate prepatterned with wells to define five of six of the walls constructing each display cell. The dyes (particles and liquid) are then ink-jet printed into each well, a top electrode is patterned, and the wells are encapsulated. The printing steps described here can all be performed via ink-jet printing, but not the patterning of the ITO substrate. The resulting product, sometimes called ‘‘electronic paper,’’ [12] has been advertised as ‘‘all printed’’ by some manufacturers, and ironically includes no paper at all (but is thin and flexible, and is based on reflection just like printed paper). This is in contrast to the ‘‘paper electronics’’ manufactured by other groups, which employ paper as a substrate in the printing process. Electrowetting is the control of interfacial energy between a liquid and a solid through an applied potential. This principle has been demonstrated in a cyan–magenta–yellow (CMY) display at video speeds [9]. In this device, the voltage applied manipulates small volumes of pigmented oil, masking or revealing the white paper-like substrate. The dynamic color switching is good and exceeds the net color change reached in, for instance, liquid crystal displays. Passive- and active-matrix memory devices are another potential application for conjugated polymers. So far these devices have not been printed, but rather manufactured in manners analogous to traditional silicon transistor processing. The advantage of the polymer devices over current silicon technology is the device density that can be achieved (through a simple-cross point matrix in the case of passively addressed memory) and stackability, which is the possibility to place one or more arrays of memory cells on top of the previous layers. Although inefficient in comparison to their silicon-based counterparts, solar cells based on conducting polymers offer an extremely inexpensive alternative power source as long as they can be processed in a roll-to-roll fashion.
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Electrochemical devices such as transistors [13–15] and diodes [16] have been demonstrated using printing techniques. These devices are quite slow (0.1 to 10 Hz), but may be acceptable for use in conjunction with printed electrochromic displays—the key advantage is that the devices and the display elements are produced simultaneously using the same process steps. Similarly, passive components and resistor–capacitor circuits have been demonstrated with ink-jet printing [16]. This last example even demonstrated the use of standard staples for interconnecting stacked layers. Linko¨ping University and Acreo AB have focused on all-screen-printed electrochromic displays with electrochemical transistors as the driving circuitry and simple logic, all powered by printed batteries. From this, logic and displays are made out of the same material combination, which minimizes the number of total printing steps. In Figure 4.3, passive-and active-matrix-addressed displays (a and b) [5] and an electronic label (c) are shown. In the display, electrochemical smart pixel circuitry allows for unique addressing of each display element in the display matrix. Both the display element and the transistor in each pixel are based on the combination of patterned PEDOT:PSS films and common electrolyte layers. The transistors employ the impedance change of PEDOT with oxidation state whereas the display elements use the electrochromic properties of PEDOT. The electronic label (Figure 4.3c) contains a printed battery, touch buttons, electrochromic display elements, and transistors to provide simple logic functions. In total, fewer than 10 printing steps are required to achieve a fully integrated electronic label. Sensors, demonstrated in a simple form by Nilsson et al. [14], are another potential application of printed organic electronics. Nilsson’s humidity sensor, based on the ionic conductivity change in Nafion with water content, is a simple system for converting ionic information to electronic information using electrochemistry based on paper coated with PEDOT doped with PSS (PEDOT:PSS). In general, more complex electrochemical sensors can be manufactured by printing. A likely candidate is the PEDOT:PSSbased glucose sensor presented by Zhu et al. [18]. Organic electronic materials are already used in commercial quality-control products (the electronic nose) and solid-state gas sensors. A conjugated polymer sensor application with a huge potential is biological detection. Since polymers are chemically synthesized, they can be tailored by adding a multitude of functions through side chains, or sometimes in the polymer backbone itself. At an elementary level, such customization changes the color of electrochromic materials [4]. Much more complex functionality can be built in, such as protein or DNA-fragment recognizing groups. These groups selectively bond with a single type of protein. The coupling can cause a rearrangement of the polymer within the film (a phase transfer), which changes the electronic conductivity.
4.1.2 Market The various markets for printed electronics must be analyzed separately. In many areas, devices based on organic materials are in direct competition with crystalline or polycrystalline silicon-based devices. Although silicon devices typically offer a significant performance advantage (as far as speed, logic complexity, etc.), printed electronics often win from a cost perspective. The flexibility of the substrate is also important for applications such as consumer product packaging. IDTechEx predicts that the market for ‘‘smart labels’’ (RFID tags for consumer packaging) will grow to over $7 billion by 2008, and over $20 billion by 2015, including all kinds of electronic devices, silicon and organic alike [19]. The challenge for organic electronics is to bring printed RFID devices to market before the silicon industry finds a way to price the new materials right out of the market.
4.2
Substrates for Printed Organic Electronics
Plastic and pulp products (paper) are two obvious choices for flexible substrates. Both are relatively inexpensive and readily available commercially. The first obvious difference between them from a product point of view is transparency. Paper is usually opaque, where plastic can be clear. The latter is often useful for displays, mirrors, and other optical components.
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FIGURE 4.3 (See color insert following page 8-22.) Printed electronics prototypes from Linko¨ping University and Acreo AB, Norrko¨ping, Sweden. (a) A seven-segment all-printed paper display. (Courtesy of Acreo AB. With permission.) (b) An active-matrix addressed plastic display. (From Andersson, P., Nilsson, D., Svensson, P.-O., Chen, M., Malmstrom, A., Remonen, T., Kugler, T., and Berggren, M., Adv. Mater., 14, 1460, 2002. With permission.)
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FIGURE 4.3 (continued) (c) An electronic label including battery, push buttons, transistor logic and display elements. (Photograph courtesy of Acreo AB. With permission.)
The unfinished surface of pulp paper often absorbs inks, which can be very useful for printing color inks, but often a challenge when producing electronics since absorbed material can be disconnected from the rest of the film. Thus, more material must be deposited if a contiguous film on top of the paper is desired. It should be noted that absorbed material typically adheres extremely well to the substrate. It can be easier to produce electronic components on paper with a thin plastic (polyethylene or similar) coating, such as that found on many photographic papers. The paper can also be finished by other means, such as impregnation or lamination, but the desired end result is a nearly impermeable surface. In this case, the printing surface is effectively the same regardless of whether the bulk of the substrate is paper or plastic. Unfortunately, adhesion between the printed film and the plastic can be more of a challenge than on unfinished paper. Surfactants and UV curing the ink (described in Section 4.5) can help in this case. LEDs require both electronic conductivity and optical transparency in one of the electrodes used to drive the device. The classic transparent electrode material is ITO, used long before conjugated polymers were considered for electronic devices. However, ITO is not printable (with the exception of a few recently developed commercial inks with relatively low conductivity), and producing ITO on flexible substrates is one of the largest costs involved in producing LEDs and similar devices. An example of ITO on flexible plastic (as used by Add-Vision Inc.) is found in Figure 4.4a. As a consequence, thin films such as PEDOT highly doped with PSS (Orgacon foil from AGFAGevaert) are considered alternatives. Although typically conductive (300 S=cm) and quite transparent (Orgacon has a slight blue color depending on the thickness and oxidation state), the thin films have not yet reached the conductivity of ITO. A picture of a prototype roll of precoated PEDOT:PSS on paper can be found in Figure 4.4b. One of the biggest advantages of separating the film manufacturing from the patterning processes is increased flexibility and control over film properties. Printing techniques typically create films with varying thickness, polymer concentration, electronic conductivity, and color, especially at the edges of features. In some cases, this condition is acceptable, but in many cases, the quality and uniformity of the film are crucial. Focusing on the quality of the film during a coating step, and later patterning the features from the film with a subtractive technique, offers more flexibility and control during film
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creation and therefore typically results in better films. Thus, the conducting film preparation step is a subject all in itself.
4.3
Coating Processes
Coating is simply the process of creating a film on a substrate, without patterning it. In other words, the film thickness is the only dimension of concern. Separating the coating and patterning steps typically simplifies the process and adds flexibility, but this flexibility comes at the expense of more processing steps. Even when considered by itself, a coating process can be very complicated. The film thickness is typically critical to the performance of the device to be manufactured. However, the conductivity of the film also depends on the concentration and structure (e.g., crystallinity) of the polymer in the film, which can also depend on the deposition rate, temperature, solvent removal rate, and a host of other parameters. Aside from the coating techniques presented here, flexo, screen, and offset printing can also be used to manufacture uniform films in roll-to-roll processes. These techniques are described in Section 4.4.
4.3.1
Spin-Coating
The coating process most familiar to those in the silicon semiconductor industry is spin-coating. Having been used to spread photoresist on wafers for over 30 years, the technique is very mature and well understood. Typically, solution is poured onto the center of a horizontal wafer, the wafer is spun around its axis of symmetry slowly to spread the solution over the whole wafer, and then the wafer is spun at high speed to remove excess material until a relatively uniform film of the desired thickness is obtained. The wafer is then baked to drive excess solvent out of the film. The process leaves a thick ring of material around the edge of the wafer (which is usually scraped-off with a blade), and the thickness of the film is typically slightly thicker in the center of the wafer than at its edge. Spinning time, angular velocity, and temperature all affect the process. The environment (especially the partial pressure of solvent and airflow rate), interaction between the solution and substrate (wetting) and the solution viscosity and concentration are also very influential.
FIGURE 4.4 Flexible conducting substrates. (a) ITO on plastic foil as used by Add-Vision. (Photo provided by Add-Vision Inc., Scotts Valley, CA. With permission.)
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FIGURE 4.4 (Continued) electrochromic displays.
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(b) Thick (400 nm) PEDOT:PSS coated on paper from AGFA-Gevaert for use in
Coating solutions of electroactive materials is very similar to coating solutions of photoresist. Thus, for small substrates, spin-coating is still widely used. Unfortunately, the process requires considerably more material than is actually used in the film. On large substrates, greater than 99% of the solution from which the film is to be cast is thrown from the substrate during coating. Recovery systems that catch this solution and reprocess it help in this situation. Circular substrates are the easiest to spin-coat, but in general most shapes that can be represented by a single-valued function r(u) in polar coordinates can be coated in this manner. The process is not applicable in roll-to-roll processing.
4.3.2 Linear Coating Processes For spreading prefabricated solutions in roll-to-roll processes, spreader bars and doctor blades are very common. The former is a thin rod (a few millimeter in diameter) on which a thin wire with a fixed
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diameter has been wound to uniformly cover the rod with one diameter of the wire. When the rod is pressed to a flat substrate and moved to push a puddle of polymer solution, the gaps between the windings allow solution to pass through, under the bar. The solution then spreads out into a uniform film. The speed at which the rod is moved, the wetting and rheologic properties of the solution and the size of the winding used all influence the resulting film thickness and quality. A doctor blade is simply a smooth bar or blade which is held a fixed distance from the substrate. The blade is then moved relative to the substrate and film precursor to spread the solution. As with a spreader bar, the velocity, rheology of the solution, interfacial energy between blade, solution, and substrate are important process parameters. Clearly the distance between the substrate and the doctor blade is important as well. In general, spreader bars are used for low-viscosity solutions and typically form thin films. Thicker films and viscous solutions call for doctor blades. Of course, in roll-to-roll manufacturing, the bar or blade can be held in place whereas the substrate (web) runs underneath.
4.3.3
Coating with Flow Processes
Incredibly complicated coating processes exist for difficult-to-manufacture films, often developed by companies making photographic films and papers for use in manufacturing their products. In many cases, solutions cascade over a weir in a very controlled fashion covering the width of the substrate. It is not unusual for multiple layers to be formed in the same process, and the solutions are often reacting (e.g., polymerizing) immediately before, during, and sometimes after, the coating event.
4.4
Patterning Electroactive Materials
Regardless of the patterning technique employed, many electronic devices place requirements on the placement (registration) and size of the areas patterned. For example, a multiple layer field-effect transistor structure that could be patterned via printing is shown in Figure 4.5. The thickness of the semiconducting layer and the length of the transistor channel are critical to the performance of the transistor, particularly the on–off ratio and driving voltages. However, the overlap d between the gate and the source and drain electrodes is effectively an unnecessary capacitor that adversely affects the switching speed of the device. This undesired overlap could mean the difference between the ability to make a small (4 by 4) or large (1000 by 1000) array of such devices in, for example, an OLED display device. The capabilities that determine the above-mentioned overlap are the degree of control over the lateral size of the printed source, drain, and gate, and the ability to register (align) the printed gate to the source and drain that were printed in a previous operation. This overlap and the registration requirement disappear if a lateral device (with only one conducting layer) can be used instead.
Gate electrode d CC C Source contact
Gate dielectric layer C
Semiconducting polymer Drain contact Carrying substrate
FIGURE 4.5 The organic field-effect transistor structure, a top-gate configuration. Addressing the gate electrode induces accumulation of charges along the channel, inside the semiconducting polymer. The undesired overlap (d) of the gate and the drain and source contacts causes the parasitic capacitance (C) that slows the transient characteristics of the transistor.
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4.4.1 Additive Patterning Additively printing materials onto substrates is the usual image that the words ‘‘printing press’’ conjure in our minds. Not only have the basic processes and equipment been available for hundreds of years, the colors and images produced surround us on a daily basis. Additive processing of materials has the distinct advantage in that nearly 100% of the material can be applied to the substrate. Little material is wasted, which can be a major factor when the ‘‘ink’’ is an expensive synthetic polymer. However, additive patterning appends rheology and ‘‘curability’’ to the electronic material requirements necessary for device operation. The printing processes described below each have their own operating conditions with regards to viscosity, drying time, curing process, etc. Thus, solutions, suspensions, or blends of conducting polymers and thickeners, fillers, cross-linkers, and solvents must be tailored for each application. In general, the films created through additive patterning techniques have a lower quality (e.g., poorer uniformity) than pure coating techniques, simply because coating is a simpler, more flexible process. From a fundamental point of view, during printing, the ink must be transported from a reservoir to the substrate. The manner in which this occurs depends on the process, described below. The ink must wet the substrate and then dry (or cure) on the substrate. The final film must adhere to the substrate, and must be mechanically flexible to maintain usefulness if the final device is to be flexible. Polymeric materials are commonly used in traditional printing processes because the resulting films meet these criteria. It requires no leap of faith to extend this principle to conducting polymers. The biggest challenge lies in maintaining the electronic properties of the film after processing. 4.4.1.1 Ink-Jet Printing ‘‘Ink-jet’’ describes small drops of ink driven to the substrate, usually by an electrostatic field. The drops are formed by a rapid pressure pulse in a small, nearly enclosed chamber. When the pressure increases, a small amount of liquid is ejected from the chamber through a nozzle (the only exit). When the pressure in the chamber decreases again, that liquid is separated from the bulk of the liquid in the chamber by surface tension in a process called pinch-off. The pressure pulse is commonly generated by a piezoelectric material or by heating the liquid via current through resistive material embedded in the wall of the chamber. The ejection process usually leaves the drop with insufficient velocity to drive it to the substrate, so another force, such as an electrostatic field, is used to complete the delivery. The result is a rapid stream of successive drops, controlled electronically. Ink-jet printing has the distinct advantage that it is a noncontact patterning process. Only the small (0.1–100 mm) drops contact the substrate, and the drops are never in contact with the substrate and the printing equipment simultaneously. A second advantage is that it is a fully digital technique, meaning that the pattern to be printed can easily be changed on pages printed immediately after one another. There is no patterned drum or mask that must be changed to print a new pattern. This is why the ‘‘best before’’ date on many consumer packages is ink-jet printed whereas the rest of the package is produced with another technique. To ensure continuous conductivity along a printed line, conducting ink drops must be printed in an overlapping fashion. Often, creating a single line or film with acceptable conductivity requires more than one layer of ink. The additional layers can be printed either before or after the previous layers have dried. Producing a continuous conducting line or film by printing overlapping drops is complicated by the evaporative drying usually used to cure the ink. Assuming that the substrate is nonabsorbing (coated with a thin film of polyethylene terephthalate (PET), for example), the receding contact line is often fixed. This means that the drop cannot shrink as the solvent is removed, resulting in a net transport of solvent (and solute) toward the perimeter of the drop [20]. The film left behind is therefore typically 10 to 1000 times thicker at the edge than in the middle. Partially drying the drop on its way to the substrate by removing some or most of the solvent can reduce this effect. When the drops produced by the ink-jet are less than about 100 mm, surface energy dominates their motion and interaction with the substrate, meaning that they do not ‘‘splash’’ or scatter material where it is not desired.
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Ink-jet printing is slow relative to techniques such as screen, flexo, and offset printing. Performance is improved by using a series of print heads in parallel, often attached to the same positioner, printing a band of the output every time the positioner sweeps across the substrate. The different heads can, of course, contain different materials. All-organic field-effect transistors made using direct ink-jet printing were reported in 2000 [21]. Complete transistor circuits were achieved by utilizing an ink-jet-printing process to make vias in insulating layers [22]. The fundamental transistor characteristics depend primarily on the drain-source channel length together with the gate and gate dielectric properties.
4.4.1.1.1
Material Requirements
One of the reasons that ink-jet printing of organic electronic materials has advanced so quickly is that the viscosity of the ink used in the process is in the millipascal second range, a reasonable value to achieve by dissolving a conjugated polymer in a solvent. The higher viscosities used in screen and offset printing are much harder to achieve without adding fillers that dilute the electroactivity of the resulting film. There are several types of ink-jet printing (e.g., thermal or piezo drop-on-demand, continuous drop formation). However, in general, viscosities between 1 and 20 mPa s produce films less than 0.5 mm thick in thermal-curing processes [23]. More viscous ink in the range of 10 to 30 mPa s can be used to produce 12 to 18 mm thick films if UV curing is employed. 4.4.1.2 Flexo Printing Flexo printing is similar to the traditional rubber-stamp techniques that have been in use for hundreds of years. The primary difference is that the flexible stamp (cliche´) in a flexo process is wrapped around a drum to be used in a continuous rotary process. In order to control the reinking of the cliche´, an anilox roller, a metal or ceramic roller with small indentations (wells) of uniform size and shape, transfers ink from a fountain to the cliche´, as shown in Figure 4.6 through Figure 4.8. Different rollers can be purchased to control the amount of ink that is transferred to the cliche´. A doctor blade removes excess ink from the anilox roller before it contacts the cliche´. Unfortunately, the contact between the soft cliche´ and the substrate can damage the substrate if the cliche´ is too hard, or wear out the cliche´ if it is too soft. Cliche´ deformation effectively limits the size of features that can be patterned whereas wear-out limits the number of devices that can be patterned without replacing the stamp. Acreo AB has demonstrated flexo printing of PEDOT:PSS films on photographic paper in a roll-toroll process. The ink base, developed by AGFA-Gevaert, was mixed with other additives to increase the viscosity of the material (see the viscosity requirements for flexo printing in Section 4.4.1.2.1).
Cliché Web direction
Anilox roller
Doctor blade
Impression cylinder
FIGURE 4.6 Flexo-printing process schematic. Ink is transferred from a well to a flexible printing cliche´ by an anilox roller, and then from the cliche´ to the substrate. The choice of anilox roller material and well size and the position of the doctor blade, which removes excess ink from the anilox roller, determine how much ink is transferred.
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FIGURE 4.7 Flexo printing of organic electronic circuits as performed by PolyIC. (Courtesy of PolyIC GmbH & Co. KG, Erlangen, Germany. With permission.)
The surface resistivity of PEDOT:PSS patterned using flexo printing is rather high, ranging from a few to 10 kV=&. However, new inks are continually developed, and improved performance is expected. PolyIC is also developing a flexo production line for RFID tags based on organic field-effect transistors. The thickness of the film that can be produced by flexo is limited both by the transfer process and the ability to cure the material. Thermal curing is effective for films up to 1 mm thick, which is sufficient for a wide range of polymer film applications. Adding photoinitiator and a cross-linking or polymerizing agent makes UV curing of films 2–3 times as thick possible at high speed [23].
4.4.1.2.1
Flexo Material Requirements
The ink used in flexo printing is relatively viscous, often 0.05–0.5 Pa s [23]. This viscosity is rather high for a conjugated polymer solution without any additives, since conjugated polymers typically have limited solubility. Nonelectronic inks often employ fillers and thickening agents to bring the viscosity within the usable range. However, these additives quickly reduce the conductivity (or electroactivity) of the resulting film, either simply through dilution or often phase separation; the latter producing films where ‘‘islands’’ of the electroactive material are electronically insulated from each other by another phase containing the filler, making the film unusable. This phase separation is not as critical in
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FIGURE 4.8 Flexo printing at Acreo AB. A flexible printing template is attached to a printing drum (a), inserted into the printer and adjusted (b),
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FIGURE 4.8 (continued) and then operated at high speed (c). (Images courtesy of Acreo AB, Norrko¨ping, Sweden. With permission.)
traditional printing, so the tricks used to adjust the fluid rheology in traditional printing are often inappropriate for electronic applications. For example, the ‘‘clay’’ often added to inks to increase viscosity may cause serious flocculation in the presence of additional ions (commonly used in PEDOT:PSS, for example). Similarly, the ions from the clay may interfere with the intended use of the polymer in the film (allowing electrochemistry, for example). Care must be taken to balance the electroactivity of the film with the materials added for printability. In flexo printing, ink must be picked up by an anilox roller, transferred to a cliche´ and then finally to the paper. Each step requires the ink to wet a new surface. Thus, the choice of materials, including solvent and possibly surfactants, is critical. As with other additives, surfactants can significantly decrease the electroactivity of the printed film if not chosen carefully and used in moderation. In the printing industry, relatively polar solvents such as ethyl acetate, alcohols, and water are most commonly used in flexo printing. These solvents are compatible with the soft rubber–like cliche´. More aggressive solvents like acetone and chloroform may shorten the cliche´’s lifetime. 4.4.1.3 Offset Printing In an offset-printing process, the printing plate has hydrophobic and hydrophilic areas instead of raised bumps. Both water and hydrophobic ink are coated onto the plate via complex transfer systems, shown only schematically in Figure 4.9. The water and hydrophobic ink displace each other so that the ink stays on the hydrophobic areas of the printing plate, thus defining the pattern to be printed. The plate-wetting process is incredibly complex. Successful transfer of the ink from the roller to the printing plate requires precise tuning of the roller and plate surface energies. The hydrophobic and hydrophilic areas of the printing plate must have a large contrast in water contact angle in order for the
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Water-transfer system Offset plate Anilox roller
Web direction
Doctor blade
Impression cylinder
FIGURE 4.9 Offset printing schematic. The oil-based ink coats only the hydrophobic areas of the printing plate. In the water-offset technique shown here, water is used to further enhance the resolution of the process. Ink and water are transferred from a complex assembly of cylinders that guide the oil ink and water to the plate in the proper manner.
ink and water separation to take place before contacting the substrate. Fortunately, at small length scales, surface tension and interfacial energy are powerful. Offset has grown in popularity because it is fast, allows for high-resolution multi-ink (multicolor) printing, and the printing plates can be manufactured in a matter of minutes (through a photolithographic process) with relatively inexpensive equipment that can be operated easily at the printing location.
4.4.1.3.1 Material Requirements Offset printing generally employs rather viscous ink, between 40 and 100 Pa s, which is hard to achieve with a polymer solution without adding fillers or other materials as described in Section 4.4.1.2.1. The resulting film is typically about 1 mm thick [23]. Liquids (inks) usually wet surfaces with which they are chemically similar. Polar liquids, those having a strong chemical dipole such as water, typically wet polar surfaces, such as glass, well. Similarly, nonpolar liquids, n-hexanes for example, wet nonpolar materials, such as Teflon, well. Unmatched materials do not wet each other; polar liquids typically have a high contact angle on nonpolar surfaces and nonpolar liquids have a higher contact angle on polar surfaces. Inks containing ions are usually extremely polar. A surfactant is a molecule, which has both a polar end and a nonpolar end. A surfactant finds its way to the interface between materials (or phases), where only small amounts (less than 2%) should be necessary to significantly lower the energy of the interface between otherwise incompatible materials. The right surfactant at the right concentration can often solve an adhesion or wetting problem. The printing industry has developed good inks and plate-manufacturing techniques for offset printing. The challenge for printed electronics is to make the ink rheologically acceptable and compatible with the ink-transfer process while maintaining the electroactivity of the final printed film. As with other techniques, the fillers, surfactants, thickeners, etc. used by the printing industry to make pigment inks compatible with offset printing dilute and hinder the performance of most printed electronics. 4.4.1.4 Gravure Printing Gravure printing is the inverse of the flexo-printing process. Ink is transferred from an anilox or other roller to a patterned plate, but the plate is designed to hold ink in the grooves between elevated lines and areas instead of on the highest points like in flexo printing. The ink is scraped from the elevated regions on the printing plate by a doctor blade before the plate contacts the web, as illustrated in Figure 4.10. The gravure process is typically chosen because of the printing speeds that can be achieved. To transfer ink at high speed, extremely low viscosity is required. Solvents such as toluene, xylene, and alcohols are often used, sometimes in conjunction with water. Fortunately, many organic polymers are soluble in toluene and xylene. The necessary viscosity range is achievable for pure polymers in solution, unlike the
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Rotogravure plate Anilox roller
Web direction
Doctor blade
Impression cylinder
FIGURE 4.10 Rotogravure printing schematic. Almost a ‘‘negative’’ of the flexo process, ink coats all of the printing plate, but is removed by a doctor blade from the elevated areas before contacting the substrate. Thus, only the valley-like wells on the printing plate transfer ink.
thicker inks required for screen and flexo printing. Removing the solvent quickly afterward, or curing the material through a cross-linking process at high speed, can be a challenge.
4.4.1.4.1
Material Requirements
As with other techniques, the material requirements for rotogravure printing depend on the curing process. For drying (solvent removal), a viscosity of between 0.01 and 0.2 mPa s is appropriate, resulting in a film about 1 mm thick [23], which is more than sufficient for many printed electronics applications. UV curing the film allows for thicker films, up to about 8 mm thick. 4.4.1.5 Screen Printing Screen printing is used outside of the electronics world for printing on all kinds of substrates, including fabrics (T-shirts) because it is rather inexpensive and highly flexible. Creating a mask (screen) can usually be done at the printing site in a relatively short time, making short printing runs of a few dozen printed products affordable. Such simple techniques rarely result in screens that can be used thousands of times. More durable masks are more expensive, and sometimes must be created off-site by companies specializing in such manufacture. Screen printing allows relatively thick films to be patterned, but lacks the lateral resolution of the other techniques presented. It seems likely that low-resolution OLED [24] and electrochromic matrix displays will be produced with this technique. It may also be used to apply protective (passivation) films to seal devices in order to prevent oxygen and water from reaching the sensitive electroactive films (see Section 4.6). Screen printing takes its name from the thin mesh of wires (called a screen) used to define the pattern. A clean, empty screen is coated with a photosensitive material. The screen (filled with the photosensitive material) is then exposed to UV or visible light through a photomask, which today can be an overhead transparency patterned in a laser or ink-jet printer. The photosensitive material hardens where exposed to light, and remains soft in the shadows produced by the photomask. The screen is then washed (developed) in a solvent that will remove the soft, but not the hardened, photosensitive material. The result is an ink mask with open areas where the screen will allow ink to pass through to the substrate, and masked areas where the hardened photosensitive material will prevent the ink from reaching the substrate. A flat-screen process as performed by Add-Vision Inc. is illustrated in Figure 4.11. A schematic of a rotoscreen process is illustrated in Figure 4.12. Even though it is not capable of the resolution that offset and flexo can produce, screen printing is a relatively easy method to implement and highly flexible as far as ink is concerned, making it an obvious first choice for low-resolution applications.
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4.4.1.5.1 Material Requirements The ink used in screen printing is typically thicker than those used in the other techniques, and sometimes benefits from complex rheology. In addition to the viscosity of the ink, the affinity the ink has for the substrate is typically very important, although it is not as complex as the wetting in offset printing, for example. Increased ink–substrate affinity (decreased contact angle) typically makes printing more reliable, simply because the ink transfers from the screen to the substrate. Ink that wets a substrate well is also more likely to produce a film that adheres to the substrate after curing. Severe shrinkage or
FIGURE 4.11 Screen printing of a polymer light-emitting device as performed by Add-Vision Inc. A pattern is created on an overhead transparency with a laser printer (a), and then transferred to the photosensitive material in a mesh. After developing, the mesh appears dark where the hardened mask material remains and light where it has been removed (b).
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FIGURE 4.11 (continued) (See color insert following page 8-22.) The screen is placed on a substrate, and an ink, in this case a blend including silver to form a counterelectrode, is drawn across the screen with a spatula (c). Ink passes through the screen only where it was ‘‘opened’’ during the developing process. The result is a pattern of ink on a flexible substrate (d).
phase changes in the added material during curing can still cause delamination, even if the adhesion is fairly strong. An extreme situation occurs when the substrate absorbs the ink, or sometimes just the solvent in the ink, as is often the case on textile fibers and paper. Although this situation is useful to dry the ink into a film, conducting material absorbed into paper or fabric will often be incongruous, resulting in inactive films. In screen printing, the contact angle between the ink and the mesh is also important. The ink should typically favor wetting the substrate over the mesh or it may not transfer reliably. To complicate matters
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Screen
Web direction
Impression cylinder
FIGURE 4.12 Rotoscreen printing schematic. Ink is pressed through a screen to the substrate in a continuous (rollto-roll) rotary process. In this case, the substrate and screen move, and the spatula remains stationary.
further, electrostatic interactions between the ink, mesh, and substrate can affect ink transfer, the quality of the resulting film, or both. This is more likely to be a problem when either the substrate or the ink has a chemical dipole or the ink contains ions. As described in Section 4.4.1.3.1, an appropriate surfactant can solve many wetting problems.
4.4.2
Subtractive Patterning
As mentioned previously, removing or deactivating material from a preformed film, leaving the desired structures behind, is advantageous in systems where simultaneous patterning and film formation is too difficult or expensive. Subtractive patterning is justified if a coating process yields superior thin-film quality. Precoated substrates are today available from a vast array of manufacturers (such as ITO on plastic foil or PEDOT:PSS on PET as in Orgacon EL from AGFA-Gevaert, both shown in Figure 4.4). Different subtractive patterning approaches are used today, such as chemical deactivation, etching, dissolution, etc. 4.4.2.1 Chemical Deactivation A chemical deactivation process is used to pattern nonactive areas in a precoated continuous film carrying specific electroactive functionality. For instance, PEDOT:PSS coatings can be deactivated using strong chemical agents such as hypochlorate. Such agents can be printed using, for instance, screen printing. The resolution of these techniques is often limited by the relationship between the diffusion of the active materials into the film, the time required for reaction, and how long it takes to remove the deactivator from the film to stop the process. As a result, the resolution that can be achieved with a simple chemical deactivation is often lower than the resolution of the printing process employed. Chemical deactivators are typically hazardous (e.g., oxidizing agents) and therefore must be handled carefully. 4.4.2.2 Electrochemical Deactivation PEDOT undergoes an overoxidation process when exposed to an anodic overpotential in an electrochemical cell in the presence of water or oxygen, known to occur at or above about 1.5 V vs. Ag=AgCl [25–28]. Originally observed as a failure mechanism in polythiophene films, this oxidation process is completely irreversible and permanently deactivates the conductivity in PEDOT and other thiophenes. The result of this deactivation is partially understood, as sulfone and carboxylic groups are observed in the remnants of the polymer after overoxidation. It appears that the conjugation in the polymer, and possibly the polymer backbone itself, is broken in the reaction. The overoxidation process is known to be very fast and can easily be incorporated into ordinary printing tools using water-based ink. The printing
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technique must allow for the formation of an electrochemical cell, where the printing tool acts as the cathode, the ink the electrolyte, and the PEDOT:PSS film as the anode [29]. Compared to simple chemical deactivation described in the previous subsection, electrochemical activation is fast, safe, and electronically controlled. Simple aqueous electrolyte (e.g., NaCl in water) is sufficient with an applied potential of greater than 5 V. High-speed processing can benefit from up to 50 V. Since the chemistry is driven by the applied electric potential, it can also be stopped instantaneously, while washing the chemical oxidant in standard chemical deactivation from the film takes time, limiting the resolution of the patterns that can be created. Figure 4.13 illustrates several techniques for patterning PEDOT:PSS via electrochemical overoxidation. Conversion of a simple plotter pen or a screen-printing station are relatively straightforward, the latter allowing print resolution of a few hundred micrometers. To demonstrate that the technique is limited only by the resolution of the printing method, 2 mm lines (the smallest mask available) were patterned in a 200 nm thick film with the aid of a photoresist mask [29].
4.4.3 Other Novel Techniques 4.4.3.1 Soft Lithography Soft lithography is a broad term that includes several techniques for patterning materials using polydimethylsiloxane (PDMS) stamps. The most applicable for this chapter is microcontact printing, since it is theoretically compatible with high-speed printing. Very similar to flexo printing, but on a much smaller scale, this process has been demonstrated to pattern electronic devices of conducting polymers down to a few microns in width [30,31]. Microcontact printing involves either depositing material on or removing material from a substrate using a rubber stamp. 4.4.3.2 Vapor-Phase Patterning of Poly(3,4-Ethylenedioxythiophene) Since conjugated polymers are chemically synthesized (or polymerized), they offer the possibility of being created in situ, rather than separately as described in the previous examples. The incorporation of processing and material manufacturing into one step increases the complexity of the process, but can yield otherwise unachievable results. The vapor-phase patterning of PEDOT [32] is one example. With this technique, an oxidizing agent is patterned onto a substrate via the lithographic technique of choice. Then, the substrate is exposed to an inert atmosphere containing EDOT monomer, which polymerizes and adheres upon reaching the oxidizer on the substrate. Where no oxidizer is patterned, the EDOT remains a monomer and is removed in a subsequent rinse. PEDOT films created with this technique have remarkably high conductivity (1000 S=cm), when compared to films manufactured with other techniques, including the coating processes described above. Patterning the oxidizing agent and creating the film in two different steps offers flexibility similar to that available by subtractively patterning a premanufactured film (described above). 4.4.3.3 Self-Organization Self-organization describes rearrangement of fluids (or even molecules) based on surface and bulk energy minimization, and can be used to enhance both the speed and resolution of printed electronics. An excellent example is the work of Wang et al. [33], where ink-jet-printed drops were printed onto a thin prepatterned hydrophobic line on the substrate. The aqueous drops, seeking the hydrophilic surface on either side of the line, split, resulting in a gap between the two film segments of around 500 nm, a length scale unachievable by ink-jet printing alone. Self-organization in the bulk is also suggested for optimizing interface-dependent devices such as LEDs and cells, and photovoltaic cells. 4.4.3.4 Embossing and Imprinting It should be mentioned that embossing (the technique used to place shiny metal foils on greeting cards and beer labels) can also be used in electronics. The technique can be used to both additively and subtractively pattern metal and polymer films.
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Counter electrode 10 V – 30 V
Electrolyte Overoxidized PEDOT:PSS Substrate
(a)
Conducting squeegee as the counterelectrode Patterned mesh
30 V
SS
PEDOT:P
Patterned mesh Gel electrolyte Cross section Overoxidized
(b)
Cross section Photoresist Overoxidized PEDOT:PSS 25 V PEDOT:PSS
Electrolyte
Counterelectrode
Patterned photoresist
(c)
FIGURE 4.13 Electrochemical overoxidation techniques for use with polythiophenes. (a) A plotter pen modified to be used in an electrochemical cell. (b) Screen-printing electrochemical cell. (c) Photoresist patterned on a PEDOT:PSS film protects areas of the film from overoxidation whereas the rest is patterned when wet by the liquid electrolyte. (Images from Tehrani, P., Remonen, T., Hennerdal, L.-O., Ha¨ll, J., Malmstro¨m, A., Leenders, L., Kugler, T., Robinson, N.D., Crispin, X., Fahlman, M., and Berggren, M., Smart Mater Struct 14, N21–N25. Copyright 2005, Institute of Physics Publishing. With permission.)
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4.4.3.5 Spray-Coating and Patterning The so-called airbrush, used to decorate everything from T-shirts to automobiles, is another option for coating or patterning conducting polymers onto substrates. The technique takes advantage of the high vapor pressure of the solvents often required to dissolve conjugated polymers. An example of reflective displays based on this technique is given by Reeves et al. [34].
4.5
Curing
Regardless of how ink is deposited on a substrate, it must be dried or cured into a plasticlike film, typically before the next processing step. Curing can be one of the most difficult steps to migrate from the laboratory to the printing press, as high-speed production means long-curing processes require a very long web pathway. The most straightforward curing technique is to simply evaporate the solvent from the solution or blend. This is often done with hot air forced along the substrate in a drying unit. As mentioned previously, perfectly reasonable drying times in a laboratory environment (say 5 min at 708C) become unwieldy in a manufacturing process. A 5 min drying step would require 100 m of web travel at a web speed of 20 m=min (a relatively slow velocity by printing standards). The 30 cm forced air-dryers common in many presses only offer about 1 s of drying time at a speed of 20 m=min. Using a low boiling point solvent (such as chloroform) for the polymeric solution can reduce the time required for drying the film. However, such solvents are usually a challenge to use safely on the industrial scale, both for environmental and personal safety reasons. Furthermore, evaporating the solvent from a film in 1 s or less may not yield optimum film properties (homogeneity, etc.). For the reasons above cross-linking polymers (either conductive or otherwise) in the ink through exposure to UV light is a popular alternative to evaporative curing. UV curing is common in many industries including the electronics industry where it is used to define submicron patterns in photoresist, which are then transferred to crystalline materials through etch processes. In roll-to-roll printing, the technique is usually used to cure patterned (or simply coated) films, rather than to pattern the films, but the concept and the basic chemistry, are similar. The ink is prepared with a photoinitiator and a crosslinking agent in addition to the starting materials normally required (solvent, conducting polymer, etc.). The photoinitiator is a light-sensitive material that forms a radical (or possibly gives up a proton) when exposed to light. This radical or proton acts as a catalyst for the cross-linking reaction, in which the cross-linking agent covalently connects neighboring polymer chains. The cross-linked film is typically a plastic adhered to the substrate. Sometimes a monomer is used with a photoinitiated polymerization agent instead of cross-linking. The end result is very similar. Theoretically, the conducting polymer itself can be polymerized from a monomer in this fashion after patterning, but such polymerizations are typically too difficult to be performed in situ. Besides dramatically decreasing the curing time, a second advantage of printing UV-curable inks is that they typically do not cling to and dry on other pieces of equipment such as the rollers, nozzles, and screens described with the various processes above. This makes cleaning and maintaining the equipment much easier. Other print industry curing techniques, such as electron-beam curing, are often too aggressive for conducting polymers to survive. Unfortunately, the electroactivity of conducting films decreases with the addition of other materials. Finding a cross-linker (or polymerizer) that does not affect the conjugated polymer can be a challenge, and even if one is found, it will decrease the conductivity of the film simply by reducing the volume fraction of the conducting material in the film. Thus, curable conducting polymer inks have been demonstrated, but are typically much less conductive than thermally cured or premanufactured films from the original polymer solution or suspension (such as Bayer’s Baytron P or AGFA-Gevaert’s Orgacon). If the substrate absorbs the solvent used in printing, as is the case with many types of paper, then the ‘‘curing’’ of the film may not require any extra help, or thermal curing may be sufficient. However, as mentioned previously, films made by absorbing the solvent (and typically some of the polymer) into the paper typically have limited electroactivity.
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Encapsulation
A large fraction of the conjugated polymer devices manufactured commercially or in research laboratories today will not function in an environment with water or oxygen, and thus must be protected from ambient atmosphere. This is an acute challenge with OLEDs, which currently seem to be manufacturable, but result in components with insufficient life span. There are, of course, multiple approaches, including varnishes and lamination of plastic foils (tape) over the device. Varnishes are used in the printing industry to protect printed surfaces and offer a glossy finish to the printed page. They are used much like other inks, and can be cured by removal of the solvent and exposure to UV light. The additional requirement of a water and oxygen barrier poses a significant challenge. Laminating a plastic foil or other barrier on top of the component to be protected is another common technique in the printing industry. By laminating a prefabricated foil, thicker and more complex protective layers can be used. These layers may offer filters to affect the color of an LED, for example. Unfortunately, organic materials are much more permeable to moisture and oxygen than metals and ceramics, for example. This is why there is a thin aluminum layer in the paper-based containers for many dairy and juice products. As mentioned previously, this is particularly important in LEDs and n-doped transistors, which are more sensitive to oxygen and water than electrochromic displays, for example. There are commercial solutions for manufacturing such a seal based on ceramics, for example, but many require vacuum processing. While compatible with roll-to-roll manufacturing, such encapsulation techniques can quickly dominate the cost of producing printed electronics.
4.7
Conclusion
Although very immature when compared to traditional printing or silicon semiconductor manufacturing, printed polymer electronics are quickly approaching commercial production, with many expecting flexible substrate products on the market as early as 2006 or 2007. The true test for conducting polymer devices will not only be the manufacturing processes and low-cost opportunities they enable, but also the new and unique functions that can be built into the polymers themselves, which will allow applications where silicon and other crystalline materials are inappropriate. The techniques described here will likely first be tested as conjugated polymers are used to replace more expensive devices, such as Si RFID tags, with less expensive printed counterparts. However, optimistic researchers will quickly point out that the applications too far removed from the ‘‘conventional’’ electronics, on which we base our experience, will define the real impact of printed polymer electronics.
Acknowledgments The authors would like to thank Luc Leenders and AGFA-Gevaert for material regarding PEDOT:PSS films and printable inks, Melissa Kreger and Add-Vision Inc. for the images of screen printing of lightemitting devices and the image of the flexible ITO on plastic substrate, Ian Parker and DuPont OLED Displays. For the image of their LED display, Wolfgang Clemens and PolyIC GmbH & Co. KG for the image of their flexo-printing process, Acreo AB for the use of their data and images of the press in the Norrko¨ping electronic printing house, and the editors for inviting us to write this chapter. The Organic Electronics group at Linko¨ping University in Norrko¨ping is part of the COE@COIN collaboration sponsored by the SSF.
References 1. Berggren, M., O. Inganas, G. Gustafsson, M.R. Andersson, T. Hjertberg, and O. Wennerstrom. 1995. Controlling color by voltage in polymer light-emitting-diodes. Synth Met 71 (1–3):2185–2186. 2. PolyIC GmbH & Co. KG. http:==www.polyic.com 3. Krumm, J., E. Eckert, W.H. Glauert, A. Ullmann, W. Fix, and W. Clemens. 2004. A polymer transistor circuit using PDHTT. Electron Dev Lett IEEE 25 (6):399–401.
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4. Groenendaal, L.B., F. Jonas, D. Freitag, H. Pielartzik, and J.R. Reynolds. 2002. Poly(3,4-ethylenedioxythiophene) and its derivatives: Past, present, and future. Adv Mater 12 (7):481–494. 5. Andersson, P., D. Nilsson, P.-O. Svensson, M. Chen, A. Malmstrom, T. Remonen, T. Kugler, and M. Berggren. 2002. Active matrix displays based on all-organic electrochemical smart pixels printed on paper. Adv Mater 14 (20):1460–1464. 6. Reeves, B.D., B.C. Thompson, K.A. Abboud, B.E. Smart, and J.R. Reynolds. 2002. Dual cathodically and anodically coloring electrochromic polymer based on a spiro bipropylenedioxythiophene [(poly(spiroBiProDOT)]. Adv Mater 14 (10):717–719. 7. Pei, Q.B., G. Zuccarello, M. Ahlskog, and O. Ingana¨s. 1994. Electrochromic and highly stable poly(3,4-ethylenedioxythiophene) switches between opaque blue–black and transparent sky blue. Polymer 35 (7):1347–1351. 8. Comiskey, B., J.D. Albert, H. Yoshizawa, and J. Jacobson. 1998. An electrophoretic ink for all-printed reflective electronic displays. Nature 394 (6690):253–255. 9. Hayes, R.A., and B.J. Feenstra. 2003. Video-speed electronic paper based on electrowetting. Nature 425:383–385. 10. Ko, H.C., M. Kang, B. Moon, and H. Lee. 2004. Enhancement of electrochromic contrast of poly(3,4-ethylenedioxythiophene) by incorporating a pendent viologen. Adv Mater 16 (19):1712–1716. 11. Acreo AB, http:==www.acreo.se 12. Rogers, J.A., and Z. Bao. 2002. Printed plastic electronics and paperlike displays. Adv Mater 40 (20):3327–3334. 13. Nilsson, D., M. Chen, T. Kugler, T. Remonen, M. Armgarth, and M. Berggren. 2002. Bi-stable and dynamic current modulation in electrochemical organic transistors. Adv Mater 14 (4):51–54. 14. D. Nilsson, T. Kugler, P.-O. Svensson and M. Berggren, Actuators B: Chem, 2002, 86, 193–197. 15. Nilsson, D., N.D. Robinson, M. Berggren, and R. Forchheimer. 2005. Electrochemical logic circuits. Adv Mater 17 (3):353–358. 16. Chen, M., D. Nilsson, T. Kugler, M. Berggren, and T. Remonen. 2002. Electric current rectification by an all-organic electrochemical device. Appl Phys Lett 81 (11):2011–2013. 17. Chen, B., T. Cui, Y. Liu, and K. Varahramyan. 2003. All-polymer RC filter circuits fabricated with inkjet printing technology. Solid-State Electron 47 (5):841–847. 18. Zhu, Z.-T., J.T. Mabeck, C. Zhu, N.C. Cady, C.A. Batt, and G.G. Malliaras. 2004. A simple poly(3,4ethylene dioxythiophene)=poly(styrene sulfonic acid) transistor for glucose sensing at neutral pH. Chem Commun (online). 19. ID Tech Ex Ltd, I. 2005. RFID market to reach $7.26Bn in 2008. 20. Deegan, R.D., O. Bakajin, T.F. Dupont, G. Huber, S.R. Nagel, and T.A. Witten. 2000. Contact line deposits in an evaporating drop. Phys Rev E Stat Phys Plasmas Fluids 62 (1B):756–765. 21. Sirringhaus, H., T. Kawase, R.H. Friend, T. Shimoda, M. Inbasekaran, W. Wu, and W.P. Woo. 2000. High-resolution inkjet printing of all-polymer transistor circuits. Science 290:2123–2126. 22. Kawase, T., H. Sirringhaus, R.H. Friend, and T. Shimoda. 2001. Inkjet printed via-hole interconnections and resistors for all-polymer transistor circuits. Adv Mater 13 (21):1601–1605. 23. Kipphan, H. 2001. Handbook of print media: Technologies and production methods. Berlin: Springer. 24. Pardo, D.A., G.E. Jabbour, and N. Peyghambarian. 2000. Application of screen printing in the fabrication of organic light-emitting devices. Adv Mater 12 (17):1249–1252. 25. Ahonen, H.J., J. Lukkari, and J. Kankare. 2000. n- and p-Doped poly(3,4-ethylenedioxythiophene): Two electronically conducting states of the polymer. Macromolecules 33 (18):6787–6793. 26. Barsch, U., and F. Beck. 1996. Anodic overoxidation of polythiophenes in wet acetonitrile electrolytes. Electrochim Acta 41 (11–12):1761–1771. 27. Tang, H., L. Zhu, Y. Harima, and K. Yamashita. 2000. Chronocoulometric determination of doping levels of polythiophenes: Influences of overoxidation and capacitive processes. Synth Met 110 (2):105–113.
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28. Tsakova, V., S. Winkels, and J.W. Schultze. 2001. Anodic polymerization of 3,4-ethylenedioxythiophene from aqueous microemulsions. Electrochim Acta 46 (5):759–768. 29. Tehrani, P., N.D. Robinson, T. Kugler, T. Remonen, L.-O. Hennerdal, J. Ha¨ll, A. Malmstro¨m, L. Leenders, and M. Berggren. 2005. Patterning polythiophene films using electrochemical overoxidation. Smart Mater Struct 14:N21–N25. 30. Xia, Y.N., and G.M. Whitesides. 1998. Soft lithography. Annu Rev Mater Sci 28:153–184. 31. Michel, B., A. Bernard, A. Bietsch, E. Delamarche, M. Geissler, D. Juncker, H. Kind, J.P. Renault, H. Rothuizen, H. Schmid, P. Schmidt Winkel, R. Stutz, and H. Wolf. 2001. Printing meets lithography: Soft approaches to high-resolution patterning. IBM J Res Dev 45 (6):870. 32. Winther-Jensen, B., and K. West. 2004. Vapor-phase polymerization of 3,4-ethylenedioxythiophene: A route to highly conducting polymer surface layers. Macromolecules 37 (12):4538–4543. 33. Wang, J.Z., Z.H. Zheng, H.W. Li, W.T.S. Huck, and H. Sirringhaus. 2004. Dewetting of conducting polymer inkjet droplets on patterned surfaces. Nat Mater 3:171–176. 34. Reeves, B.D., C.R.G. Grenier, A.A. Argun, A. Cirpan, T.D. McCarley, and J.R. Reynolds. 2004. Spray coatable electrochromic dioxythiophene polymers with high coloration efficiencies. Macromolecules 37:7559–7569.
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II Applications and Devices Based on Conjugated Polymers
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5 Polymers for Use in Polymeric LightEmitting Diodes: Structure–Property Relationships 5.1
Poly(phenylenevinylenes).................................................... 5-6 Background Disorder
5.2
.
Copolymers
.
Morphology and Energetic
One-Dimensional Small PPV-Based Molecular Model Compounds for the Study of Polymers .......................... 5-16 Effect of Substituent Addition: Changes in Morphology Effect of Substituent Placement: Changes in Optical Properties
5.3
.
Two-Dimensional Small PPV-Based Molecules: Effects on Charge Delocalization ..................................... 5-18 Intrinsic Two-Dimensional Charge Delocalization in PPV Derivative . PPV-Based Dendrimers . Two-Dimensional Conjugated Molecules with Four Arms and a Conjugated Core
Hermona Christian-Pandya, Subramanian Vaidyanathan, and Mary Galvin
5.4 5.5 5.6
Polyfluorenes...................................................................... 5-22 Phosphorescence in PLEDs .............................................. 5-24 Summary ............................................................................ 5-26
In 1977, with the report of high conductivity in dope polyacetylene [1,2], interest in the field of conductive polymers surged. While this research eventually led to Professors Heeger, MacDiarmid, and Shirakawa receiving the Noble Prize in Chemistry, commercialization of conductive polymers was impeded by poor stability, and research activity decreased until a 1990 Nature paper by Friend and colleagues was published [3], which reported on a light-emitting diode with poly(p-phenylenevinylene) (PPV) as the light-emitting layer. The efficiency of this device was low at 0.01% internal quantum efficiency [3]; the PPV used as the electroluminescent (EL) layer emitted in the green and was synthesized via a precursor route, since it was not solution-processable. Since then there have been considerable advances in polymer design and devices. Phillips introduced an electric shaver with a polymeric light-emitting diode (PLED) display in 2002 and reported efficiencies of 14 Cd=A in the green
5-3
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with lifetimes in excess of 104 h [4]. Other corporations are considering PLEDs for commercialization in displays, although at this time evaporated small molecule organic light-emitting diodes (SMOLEDs) have captured more of the market. Some of the advantages that PLEDs have over liquid crystal displays are that PLEDs do not require a backlight and need fewer or no filters. They have faster response times, an advantage for video, a 1808 viewing angle, may be printed [5,6], and can be made on flexible substrates [7], although encapsulating a PLED made on a flexible substrate to ensure long life is still a major challenge. A challenge facing the field is, however, to improve lifetime of the PLEDs, particularly for blue-emitting devices [7]. In terms of polymer design, side groups have been added to PPV to make it solution-processable and to tune the emission color. Additionally, the class of polymers utilized in PLEDs has expanded beyond substituted p-phenylenevinylenes to include homopolymers and copolymers of p-phenylenes [8–10], 2,7-fluorenes [11,12], and 2,5-thienylenes [13,14] (see Chart 5.1 for structures of polymers and abbreviations used through out this text). A quick review of the literature reveals that there has been a plethora of research on polymers for use as active layers in PLEDs. It is now possible to tune the emission of the polymer from blue to green to red, to engineer the HOMO and LUMO levels and, to vary the architecture of the chain from linear to dendritic.
*
Poly(phenylenevinylene) (PPV)
* n R
Poly(p-phenylenes) (PP) *
*
n
R
R
*
*
n
CHART 5.1
Structures of polymers discussed in this chapter.
Polyfluorenes (PF)
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O *
MEH-PPV
* n O
OC6H13 OC6H13 CN
* C6H13O
NC C6H13O
n
*
n
O
* n
S
cyano-PPV
*
* O
*
PEDOT:PSS SO3H 3
* n
*
Poly(n-vinyl carbazole)
N
PVK
RO
OR
dialkoxyphenyl-PPVs * n
*
CHART 5.1 (continued) (continued )
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R R
N
Diamino-PPV
* n N
*
R
R
N
N O C16H33 *
*
Pyridopyrazine -PV copolymer (near IR emitter)
N n C16H33O
CHART 5.1 (continued)
With all these developments, it is beyond the scope of this chapter to review the entire field. Even a review of the synthesis of polymers for PLEDs would be too extensive for this chapter. Instead this chapter will focus on some of the structure–property relationships that govern the performance of EL polymers. A quick review of poly(phenylenevinylenes) and polyfluorenes will be discussed, since they and their copolymers have emerged as the most extensively studied polymers. Examples of research done in the Galvin group will also be included in order to highlight some structure–property relationships. For further information the reader is referred to other reviews of the field. Holmes has published excellent reviews on the chemistry and synthesis of conducting polymers [15,16]. The synthesis and properties of polyfluorenes and their use in PLEDs is covered by Wu and colleagues in a recent review [6]. For a review of many different polymer classes used in PLEDs and some of the physics beyond the properties of these materials as well as their performance in PLEDs, the reader is referred to a review by Akcelrud [17]. Finally, Schwartz has written an excellent review of how chain conformation and film morphology influence the photoluminescence (PL) and energy transfer in some light-emitting polymers [18].
5.1 5.1.1
Poly(phenylenevinylenes) Background
As mentioned earlier, PPV was the first polymer used in a PLED. The device architecture of this PLED is shown in Figure 5.1. The EL polymer, PPV in this case, was sandwiched between a transparent conductor, indium tin oxide (ITO), as the anode and aluminum as the cathode. In this configuration, when a voltage is applied to the electrodes, holes are injected into the HOMO of the PPV at the ITO– PPV interface and electrons are injected into the LUMO of the PPV at the Al–PPV interface. These
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Al
V PPV
ITO coated/Glass
FIGURE 5.1
Device configuration of first PLED.
carriers, holes and electrons, migrate through the nominally 100 nm thick PPV film and some of them form excitons, a fraction of which decay to the ground state by emitting light. The efficiency of the device is a ratio of the photons emitted, as measured by a photodiode, to the current flowing through the device. In this first device, the quantum yield was low (0.01% internal quantum efficiency) [3], in part due to defects in the polymer and in part due to poor injection of carriers from the aluminum and ITO electrodes. To understand some of the structure–property relationships important in PLEDs, it is worth discussing some of the history of how PLEDs evolved from this first device to those of the present. The PPV used in this PLED was synthesized via a Wessling [19] route, shown in Figure 5.2. The sulfonium precursor polymer is soluble in methanol, allowing one to cast very thin films (about 100 nm) of it onto ITO-coated glass using spin-coating. The precursor was converted to PPV when heated above 2008C in vacuum. The resulting PPV polymer is very crystalline and insoluble. Following this discovery that PPV could be used as the active layer in a light-emitting diode, researchers at Bell Laboratories determined that carbonyl groups were formed on the vinyl bonds of the polymer during thermal conversion [20,21]. The carbonyl groups acted as traps for electrons and so seriously limited the PL of the PPV [22] that methods of preventing their formation were sought. It was determined that if the precursor was converted to PPV in a mixture of 15% H2 in 85% N2, no detectable carbonyl groups were formed [20], and PL and EL efficiencies increased significantly [21]. While annealing in the presence of H2 limited carbonyl formation, it did not eliminate PL quenching often seen in crystalline materials. To accomplish this, the researchers at Bell Laboratories replaced the tetrahydrothiophene leaving group with an ethylxanthate group. The ethylxanthate group could be eliminated at lower temperatures, preventing carbonyl formation. The PPV formed after thermal conversion was a copolymer of cisPPV, para-PPV, and units in which the ethylxanthate group was still attached (Figure 5.3) [23]. The presence of the cis linkages prevented the polymer from crystallizing and significantly enhanced PL and EL efficiencies. An alternate way to increase efficiency was discovered by Holmes and colleagues [24]. They replaced some of the tetrahydrothiophene groups with methoxy groups, which can be eliminated
S+
Cl−
Cl−
Base
S+
* +
Cl− S
* n Heat
* * n
FIGURE 5.2
Wessling precursor route to PPV.
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O [
] [ x
y[
S
S [
z
]
FIGURE 5.3 This copolymer was prepared from a precursor in which the tetra-hydrothiophene units employed by Wessling were replaced by ethylxanthates. When heated, not all of the xanthate groups were eliminated and both cis and trans linkages were formed, as opposed to the all trans PPV prepared by Wessling. Incorporation of cis linkages prevented crystallization and enhanced PL and EL efficiencies.
via acid treatment, but not through thermal annealing. This allowed the researchers to control the conjugation length and HOMO and LUMO levels in the polymer via the elimination procedure. When the precursor polymer was only heated, the remaining methoxy groups limited conjugation in the polymer and increased EL efficiency by 30 times. These early experiments provided some useful information on structure–property relations, specifically the importance of crystallinity and conjugation length. The insoluble nature of PPV, however, made it difficult to work with and limited commercialization of PLEDs based on PPV. To make PPV soluble, side groups were affixed to the phenyl ring. The first report of a PLED made from a soluble PPV was published in 1991 [25]. The polymer used, poly[2-methoxy-5-(2-ethylhexyloxy)-1,4-phenylenevinylene] (MEH–PPV), is one of the most extensively studied polymers to date. The discovery of MEH–PPV was important not only because it proved that soluble, and therefore commercially viable, forms of PPV could be made, but also because the light emitted was red, extending the fraction of the visible spectrum that could be used and beginning the idea of making full color displays from PLEDs. Shortly after this publication, Holmes and colleagues [26] reported on the synthesis of a cyano-containing PPV, made via a Knoevenagel condensation, that contained alkoxy groups on the ring for solubility and cyano groups attached to the vinyl bonds to lower the LUMO level of the polymer. Cyano-PPV, also a red emitter, was important in two respects. The authors showed that by lowering the LUMO of the polymer, they could more efficiently inject electrons into the polymer with an aluminum cathode as opposed to using low work function cathodes such as calcium. Since this time there has been considerable research dedicated to engineering the HOMO and LUMO levels of polymers in order to align these energy levels with those of the anode and cathode, thus, improving carrier injection and the performance of the polymer in PLEDs and PVs. The second major step forward from this research was that the authors used two active layers in their PLED. They coated PPV precursor onto ITO and thermally annealed it to get an insoluble layer of PPV. They then spun the organic soluble cyano-PPV on top of the insoluble PPV. Since the HOMO level of cyano-PPV did not allow for efficient injection of holes, the use of the two-layer structure made it easier to inject both holes and electrons, yielding a device with 4% internal quantum efficiency. Presently researchers and manufacturers still use this concept, although instead of using insoluble PPV for hole injection, they use a hole injection layer (HIL) like water-soluble poly(3,4ethylenedioxythiophene)–poly(styrenesulfonate) [27,28] (PEDOT–PSS). Since these HIL layers are not soluble in organic solvents, it is possible to spin a second light-emitting layer on top of it. After the publications on cyano-PPV and MEH–PPV, researchers reported on a vast number of PPV polymers incorporating different groups that would impart solubility and in some cases change the HOMO and LUMO levels of PPV to tune the color of the emitted light. Red light could be obtained from diamino-substituted PPVs [29,30], in addition to MEH–PPV and cyano-PPV. Blue emission was reported from a PLED made from a blend of a phenyl-substituted PPV and poly(9-vinyl carbazole) [31]. Green light was emitted by dialkoxyphenyl-PPV [32] and dimethoxyoctylsilyl-PPV [33]. To enhance PL
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efficiency by separating polymer chains and reducing quenching, Wudl [34] prepared PPV with two cholestonoxy groups that emitted yellow light. Alternatively, Bao [35] substituted the phenyl ring of PPV with dendimers to enhance PL. In both cases, the EL efficiency of PLEDs made with these polymers was not high, perhaps because while enhancing PL, the increased spacing between chains hurt transport. In addition to exploring homopolymers of PPV, researchers have examined many other variables in an effort to increase their performance in PLEDs. Specifically, researchers have looked at copolymerization and chain architecture as a way to tune the properties of the materials. They also determined that energetic disorder and the morphology of the film spun from the polymer solutions had a large impact on the efficiency of the device. Each of these variables as they relate to PPV-based polymers will be discussed briefly in the following sections.
5.1.2 Copolymers Copolymerization has been widely adopted in PPV derivatives for LED applications as a tool to control defects in the polymer, achieve good solution processablity and film morphology control, tune the emission color by band gap engineering, improve transport properties by energy level engineering, or achieve high luminescence efficiency by exciton confinement. In this section, we review the progress in this area, paying particular attention to the literature over the past five years. Color tuning of emission from PPV derivatives has been achieved by the incorporation of comonomers that either interrupt conjugation and thus, confine the exciton delocalization or possess vastly different chromophores. This incorporation, in most cases, has been statistical, given the synthetic ease for such architecture. Vaeth and coworkers [36] have statistically incorporated parylene-N in the chemical vapor deposition of PPV to obtain blue emission from the polymer (Figure 5.4). Parylene disrupts conjugation in PPV confining the exciton to fewer units in the backbone, thus, shifting the blue emission. By this method, depending on the monomer delivery rates, the authors have demonstrated EL color tuning from 420 to 525 nm. Concomitant with the exciton confinement, the electrical conductivity of the polymers decreases as evidenced by significant increase in the turn-on voltages for these devices (up to 27 V) compared to pristine PPV, with a direct dependence of the voltage on parylene content. The maximum EL efficiency achieved in this copolymer is 0.05% for 60% parylene incorporation, significantly higher than similarly deposited PPV ( 0.002%). Other groups [14,37–40] have adopted a more conventional approach to color tuning, viz. copolymerizing with different chromophores such as aromatic amine segments, thienylene vinylene, spirobifluorenes, and phenyl-substituted PPVs. Becker et al. [38] report color tuning from green to orange, varying the amount of alkoxy PPV incorporated into alkoxyl-substituted 2-phenyl PPVs. With this approach, they report luminance efficiencies of ca. 15 cd=A and power efficiencies of 16 lm=W at 100 cd=m2 for two-layer devices ITO=(PANI or PEDOT)=polymer=(Ca or Yb). The authors note that the emission color depends superlinearly on the alkoxy PPV content, suggesting that a majority of the emission occurs from the sequences of alkoxy PPV in the copolymer. Reynolds [41] has designed alternating donor–acceptor copolymers that emit in the red or near infrared. Of particular interest was the alternating copolymer (AC) of the donor 1,4-(2,5-dihexadecyloxyphenylene) and the acceptor 5,8 linked 2,3-diphenylpyrido[3,4-b]pyrazine. The polymer had a band gap of 1.8 eV with emission centered at 800 nm. No efficiencies for PLEDs fabricated from this polymer were reported, but turn-on voltages were relatively high at 8V.
*
* n
m
FIGURE 5.4 By incorporating parylene units into a statistical copolymer with phenylenevinylene, Jensen was able to break conjugation and obtain blue emission.
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Recently, Mu¨ller and coworkers [40] report PPV copolymers containing fluorene units with oxetane side groups that allow photochemical crosslinking. This approach makes the polymer amenable to highresolution patterning techniques for multicolor displays. Maximum efficiencies of 7 cd=A from this green-emitting polymer have been achieved. Pixilated displays of resolution approaching a few microns have been reported with this approach. Martin et al. [42] have obtained substantial reduction in turn-on voltages without degradation in EL efficiencies by copolymerizing poly(2,3-dibutoxy-1,4-phenylene vinylene) with PPVs containing silyl side chains. Double-layer devices with Ca cathodes and employing PEDOT for hole transport yielded luminance efficiencies up to 0.72 cd=A and turn-on voltages of 4.0 V. Alam and Jenekhe [43] have investigated alternating copolymers of benzobisazole and phenylene vinylene as electron transport materials for PLEDs. With this polymer for electron transport, the authors have demonstrated reasonable improvements in luminances and quantum efficiencies with reduction in turn-on voltage in both MEH–PPV and PPV devices. Cao and colleagues have synthesized PPV-containing copolymers with 1-methoxy-4-octyloxyphenylene (MOP), and 2,1,3-benzothiadiazole (BT) units. As the concentration of BT increased, the band gap of the copolymer decreased and the emission was red-shifted [44]. An alternating copolymer of dioctyloxy phenylenevinylene and BT, when blended with [6,6]-phenylC61-butyric acid methyl ester (PCBM), was later shown to be useful in photodiodes [45]. Galvin and colleagues have recently synthesized copolymers between an oxadiazole-containing PPV derivative and alkoxy PPV (Figure 5.5) in an effort to tailor the energy levels and enhance transport [46]. All of the copolymers had 50 mol % oxadiazole–phenylenevinylene units, shown in red, and 50 mol% phenylenevinylene units, shown in blue. Since incorporation of the electron-deficient oxadiazole moiety into the PPV backbone lowers the HOMO and LUMO [47,48], it provides a tool to systematically engineer these levels in a copolymer and imparts resistance to oxidation at the vinyl bond. With an alternating copolymer, the HOMO and LUMO levels are between those of the homopolymers, which has been previously confirmed by cyclic voltammetry in conjunction with optical absorption measurements [47,48]. While the alkoxy PPV is electron injection limited in a single-layer LED with an aluminum cathode and ITO anode, OxaPPV and alternating copolymer are hole-limited to varying extents. Devices with an alternating copolymer perform significantly better than either of the devices made with the homopolymers, indicating a better charge balance even though this polymer is still hole limited. Replacing aluminum for gold cathode in this device did not impact the device performance considerably. The absorption spectrum of dilute solutions of alternating copolymer in tetrahydrofuran falls fully within the homopolymers spectra, thus, revealing a peak maximum and absorption edge midway between the homopolymers. With a statistical copolymer of 50:50 composition (SC50) between PPV and OxaPPV, the absorption spectrum in solution is broad, with good overlap with OxaPPV at higher energies, and PPV at lower energies (Figure 5.6). SC50 is not a physical blend of two homopolymers, since the OxaPPV in SC50 with C8H17 would not be soluble as a homopolymer. Note that the peak maximum of SC50 matches OxaPPV and the absorption edge matches PPV. This is best explained by the intrachain and interchain variations of oxadiazole composition in SC50, thus, there are polymer chains comprising long sequence runs of OxaPPV, long sequence runs of PPV, and a range of sequences in between, keeping the overall composition still 50:50. In this scenario, the HOMO and LUMO levels are expected to follow the oxadiazole composition and thus, vary dramatically across the chains, yielding a spread in these energy levels. Indeed, devices from SC50 show over 40% improvement in efficiencies as compared to an alternating copolymer in a single-layer configuration (Figure 5.5), attributable to decreased barriers at the injecting electrodes and improved transport across the chains. The statistical copolymers also exhibit remarkable tolerance for composition, with little change in device characteristics from SC50 to SC30 (30% PPV) as seen in Figure 5.7. With increased PPV content, however, the polymer SC70 exhibited a drop in EL efficiency. [Despite this decrease, the efficiencies with SC70 compared very well with the efficiencies of an alternating copolymer.] Flexible PLEDs were prepared with RC30, since device efficiencies with an Al cathode were good. In the a device structure of ITO=PEDOT:PSS=RC30=LiF=AL, an external quantum efficiency of 0.8%, a luminance of 4 cd=A, and a power efficiency of 1.8 lm=W, for a PLED on plastic were obtained[49].
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Polymers for Use in Polymeric Light-Emitting Diodes
Polymer
EL efficiency (ITO/Polymer/Al)
C8H17O (
]
(
y
n
0.047 %
OC8H17 PPV N N C8H17O
O
*
* n
0.05%
OC8H17 Oxa-PPV
OC8H17
OC8H17
AC
] n y
OC8H17
0.14%
SC
N N O OC12H25
C12H25O (
C12H25O (
OC8H17
C8H17O ( x (
C8H17O [ (
NN O
0.10%
)n
m
(
(
(
C8H17O
(
NN O
C8H17O
OC12H25
n
0.02%
BC
FIGURE 5.5 Structure of oxadiazole and phenylenevinylene copolymers. Note that device efficiency varies with the sequence distribution of the copolymers even though all the copolymers have an overall composition of 50% oxadiazole and 50% phenylenevinylene.
In a continued attempt at understanding how sequence distribution affected PLED performance, Vaidyanathan [46] in the Galvin group synthesized and investigated a multiblock copolymer between PPV and OxaPPV. This synthetic strategy precluded interruptions in long sequence runs of both OxaPPV and PPV. As a result, the HOMO–LUMO levels should be characterized by the pair of levels corresponding to the homopolymers. While this may improve injection characteristics at the opposite electrodes (with holes injected into PPV segments at the polymer–Al interface and electrons into OxaPPV segments at the ITO–polymer interface), recombination efficiency is expected to be severely impaired due to trapping of the opposite carriers in different segments of the chain. If there is phase separation, which is not likely with short block lengths, or isolation of the individual domains, it could provide parallel pathways for carrier transport across the electrodes, reducing the probability of recombination. Devices with BC50 show severe degradation in EL efficiency in a single-layer
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0.12 RC50 0.1
OxaPPV
Abs. Intensity (a.u.)
PPV 0.08
0.06
0.04
0.02
0 375
400
425
450
475
500
525
Wavelength (nm)
FIGURE 5.6 Note that the absorption spectra of the random copolymer SC50 (RC50 in legend) overlaps that of the Oxa-PPV homopolymer at high energies and PPV at low energy.
configuration. BC50 is among the first reports of conjugated–conjugated block architectures [50,51] and the direct donor–acceptor type of linkages in the polymer promise a unique pathway to tailor the electronic properties for applications such as polymeric photovoltaics. Hwang and coworkers [52] have copolymerized fluorenyl vinylenes with PPV (poly(FV-co-PV)s) to obtain an order of magnitude improvement in luminance efficiencies (1 cd=A) compared to poly(fluorenyl vinylenes) (PFV) (0.13 cd=A). While the authors prefer to attribute this to an increase in the HOMO level of the copolymer compared to PFV, facilitating better hole injection, it is not clear how this trades off with the impaired electron injection from an aluminum cathode in this case, since the
Polymer
Composition (% PPV)
EL eff. (%) ITO/Poly/Al
EL eff. (%) -Plastic ITO/PEDOTPSS/Poly/LiF/AL
0.8% 4 Cd/A 1.8 lm/W
SC30
30 %
0.14 %
SC50
50 %
0.14 %
SC70
70 %
0.10 %
FIGURE 5.7 The statistical polymers of oxadiazole–phenylenevinylene and phenylene are relatively insensitive to composition. The copolymer SC30 is particularly good in a PLED fabricated on a plastic substrate with PEDOT:PSS for hole injection and LiF=Al for the cathode.
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band gap remains constant. It is possible that the composition variation they have achieved in their copolymer could have a similar effect that we have suggested, improving carrier transport in their devices. Sun and coworkers [53] report a dramatic enhancement of PL quantum efficiency by introducing flexible nonconjugated spacers in PPV derivatives. With such a design, PL efficiencies over 90% have been recorded in films of the alternating block copolymer (ABC). This is a threefold increase in comparison with the homopolymer and a model oligomer. Disruption in order and limiting interchain exciton formation have been suggested as an explanation for this. Similar results have been obtained by Bradley and coworkers [54,55] who report time-resolved PL measurements, indicating reduction in excimer emission and hence an increase in PL quantum yield in MEH–PPV by incorporation of low amounts of ethylene linkages statistically in the polymer. Cacialli and coworkers [56] also report solid-state PL efficiencies of 34% and luminance efficiencies as a single-layer device between ITO and aluminum electrodes of 0.5 cd=A in poly(distrylbenzeneblock-hexa(ethylene oxide)). The authors note that in addition to confining the exciton delocalization, the oligo(ethylene oxide) spacers allow good solubility of the polymer, eliminating the need for side chains of the PPV backbone. This is particularly advantageous in maintaining the planarity of the chromophore. Karasz and coworkers have also interrupted conjugation in PPV derivatives with flexible spacers [57,58] or with meta-linkages [58] to isolate chromophores. Selective thermal elimination of PPV precursors [23,59–61] to obtain statistical copolymers has also been extensively used in interrupting conjugation to suppress aggregation.
5.1.3 Morphology and Energetic Disorder While the chemical structure of the polymer or PPV chain has a significant impact on the color of light emitted by the PLED and on the PLED efficiency, chemical structure is not the only material property that is important to the performance of the device. As mentioned in earlier sections, the LUMO level of the polymer should be low enough to allow for injection of electrons from the cathode and the HOMO level should align with the work function of the anode to afford facile hole injection. Other factors that need to be considered are energetic disorder and the morphology of the film formed via spin-coating. To explore the significance of these variables, we will consider their impact on PPV-based systems, but note that the importance of these parameters is not limited to PPV-based materials. Early in the development of PLEDs, it was noticed that different groups obtained very different device efficiencies, even for the same polymer and device configurations. In part this was due to the art involved in fabricating a PLED. In fact, the reader is cautioned about comparing efficiency results from different research groups. Slight differences in fabrication procedure can cause significant differences in efficiency, but this did not explain all of the results. Some differences were the result of variations in film morphology, specifically differences in packing between chains, caused by variations in solvent, polymer concentration or processing conditions. As noted previously, when the distance between chains is increased, PL increases, but transport does not increase, and EL efficiency depends on both properties. By changing the length of the alkoxy groups used to solubilize PPV, Vaidyanathan et al. [62] determined that there is an optimum alkyl length for improving PLED efficiency in these PPVs. This optimum length was ascribed to obtain a balance between increasing PL efficiency by separating chains and decreasing injection and transport with chain separation. Schwartz has written a comprehensive review on how morphology impacts PL and devices [18]. In his review article, Schwartz notes that interchain energy migration is faster than intrachain migration. This was shown in experiments he did on MEH–PPV [63–65] and in experiments by Lee and colleagues on polyfluorenes [66]. While this concept seems counterintuitive, the argument is that in amorphous polymer films where some segments are twisted and, therefore, not conjugated with other segments of the same chain, energy migration down a chain requires physical rotation or reorientation of chain segments to afford conjugation between these segments. This process of reorienting chain segments is slow. In Schwartz’s words, the polymer chain must ‘‘iron out the kinks’’ [18] for intrachain migration to occur. For interchain, or through space migration, to occur, the segments must be in close proximity,
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thus, once this condition is met, interchain energy migration is fast. This has important implications for the PL efficiency of polymer films, for carrier transport and injection and, therefore, for the EL efficiency. Photoluminescence from isolated polymer chains in solution can be very high, but the PL from polymer films is often considerably lower. This can be attributed to portions of the film where chains are closely packed or aggregated enough to permit the formation of interchain excitations that have very poor PL efficiency. Rapid interchain energy migration allows the energy to find these sections of aggregated chains, which are lower in energy and red-shifted with respect to the isolated chain segments [67]. The significance of these aggregated portions was also studied by Rothberg in film and in solution. By adding a nonsolvent to the solution of MEH–PPV in a good solvent, Rothberg was able to study the aggregated state with solution techniques such as nuclear magnetic resonance (NMR). Combining these studies with absorption, transient PL and steady-state PL, he developed a two species model for MEH– PPV films [68]. The two species, which coexist in the film, are isolated chain segments and aggregated species. The presence of aggregated species was shown to reduce PL by as much as two-thirds of its original value. These results began to explain variations in PLED efficiency between groups. Depending on the solvent used to spin the film, a greater or lesser degree of chain aggregation occurred in the film. In research predating Rothberg’s two-species model, Schwartz noted the existence of aggregates in MEH–PPV solutions [69] and showed that at concentrations used to spin-cast films of MEH–PPV for PLEDs, there was a significant fraction of these species. Studies involving x-ray diffraction [70], electron diffraction [71] and atomic force microscopy (AFM) [72] have also shown that there are aggregated chains in films of MEH–PPV and that there is a variation in the extent of aggregation depending on the solvent used for spinning. While aggregation seriously decreases PL, its effect on device performance is not so straightforward, since aggregation increases carrier transport. If carrier transport is low, the carriers accumulate at the polymer–electrode interface, which leads to ‘‘space-charge-limited injection’’ [73]. In short, if the film consisted of only isolated chains, there would be fewer carriers and they would move through the film more slowly. Both of these effects would decrease electroluminescence, all other factors being equal. In reality, all other factors are not equal: PL increases with isolated chains. Thus, to maximize electroluminescence efficiency, there is a trade-off to be made between increasing PL and increasing the number and mobility of carriers. This balance is highlighted by Schwartz’s research on how the efficiency of MEH– PPV PLEDs varies with the solvent from which the films were cast. Devices were made from films spun from solutions of MEH–PPV in tetrahydrofuran (THF) or chlorobenzene (CB). THF is a good solvent for MEH–PPV, thus the films had a higher fraction of isolated chains than did the films spun from CB. The PLEDs made from the THF solution had a lower current density and brightness, but the EL efficiency was higher [72]. In this case, while the isolated chains meant that the device had fewer carriers, once the carriers formed an exciton, the high PL yield of those excitons provided for a more efficient device. Where the exact balance lies will vary with system, and points to the large impact that processing can have on device performance. Yang took this concept further by relating processing-induced morphology, as influenced by solvent, concentration and spin rate to device performance [74,75]. With the exception of biological polymers such as proteins and DNA, polymers are inherently mixtures. Even in homopolymers, the difference in length between chains can be significant. Electrochemical studies have shown that in PPV systems, the energy level of the HOMO moves up and the energy level of the LUMO moves down as the conjugation length increases [76]. For conjugated polymers, the difference in chain length, coupled with differences in defect concentrations and distributions between chains, mean that the conjugation length varies between chains and chain segments. In other words, a distribution of energy levels, or energetic disorder, is present in films of conjugated polymers. Theories by Bassler [77] on transport of carriers via hopping through a material with a Gaussian distribution of site energies indicate that transport is dependent on the width of the distribution. It is, therefore, reasonable to speculate that the distribution in conjugation length, or energetic disorder, present in polymers could affect PLED efficiency. To evaluate the effect of energetic disorder on PLED efficiency, Galvin and colleagues took two approaches. The first approach employed PPV oligomers and polymers. Specifically they synthesized
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and characterized a pentamer (5PV), a nanomer (9PV), and a fairly monodisperse polymer of an octyloxy-PPV (Mn ¼ 4200, PDI ¼ 1.3) [78]. Optical studies indicated that 9PV and PPV had the same band gap (2.4 eV), whereas 5PV had a higher band gap of 3.0 eV. Thus, by mixing 9PV or PPV into 5PV, the authors could create energetic disorder in the films. PLEDs with an ITO=PV=Al configuration were fabricated from 5PV, 9PV, PPV, and blends of 0.5% PPV in 5PV or 0.1% of 9PV in 5PV. Mixing 0.5% of PPV into 5PV dropped the efficiency of the PLED by a factor of 3 as compared to the 5PV PLED. This occurred despite the fact that the efficiency of a PPV-based PLED was twice that of a 5PV-based PLED. When 0.1% of 9PV was blended with 5PV, the efficiency of the PLED dropped by a factor of 6 as compared to the 5PV PLED. The PL efficiency of the blends was the same as that of 5PV. Additionally, NMR and photophysical studies showed that the PPV was incorporated into the 5PV matrix and not excluded from it [79]. This data is a strong indication that energetic disorder, created by mixing a low band gap material into a higher band gap matrix, degraded device performance. To further probe the significance of energetic disorder, the authors fractionated an oxadiazolecontaining PPV polymer (Oxa-PPV) [78]. In this case, a nonsolvent was added to a solution of OxaPPV in a good solvent. Several fractions of the polymer, differing in molecular weight, were collected. Once again the authors fabricated PLEDs from the individual fractions and blends of 2% of a high molecular weight fraction in a matrix of low molecular weight polymer from the same synthetic batch of the material. There was no difference in PL between the fractions and the blends. There was, however, at least a 50% drop in the efficiency of a PLED made from the blend as compared to one made from the matrix. Since in both the oligomer and fractionated polymer studies, PL was unaffected by the energetic disorder created through blending, and EL was seriously degraded, it is likely that energetic disorder affects injection and transport of carriers in the blend. The importance of this parameter is highlighted in Table 5.1. All of the PPV polymers were synthesized in the Galvin laboratory. As the molecular weight decreased so did polydispersity and most likely energetic disorder. Note that there is a 100 fold difference in performance between the polymer and the lowest molecular weight PPV. The data from the Galvin group illustrates that energetic disorder can be a significant problem and that the source of the problem is not degradation in PL efficiency, but instead comes from problems with injection or transport. To obtain a deeper understanding of the influence of energetic disorder, Rothberg [80], in collaboration with Smith, extended the device model of OLEDs developed by Davids et al. [81] to include the existence of charge traps and energetic disorder in the active layer of a single-layer LED. The model allows one to predict device efficiency, quantum efficiency, the internal field, the recombination zone, and the current voltages of a device, assuming the energy levels and the carrier mobilities for the active layer and the electrodes. Figure 5.8 is useful for discussing how energetic disorder can affect a PLED. The figure indicates that the HOMO and LUMO levels are really a distribution of levels, with the solid lines representing the HOMO and LUMO levels if there was no energetic disorder. Depending on the relative positions of the work functions for the electrodes and the HOMO and LUMO levels, carrier
TABLE 5.1
PPV PPV PPV PPV
block Low Medium High
The Effect of Polydispersity on PLED Performance Mw
EL Efficiency (ITO=Polymer=Al) (%)
1750 4100 9200 125,000
0.180 0.080 0.011 0.001
The effect of polydispersity is seen in these numbers. As the molecular weight decreases, so does polydispersity. In essence, there are no chains with extremely long conjugated segments that can act as traps.
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injection could be increased or decreased by energetic disorder. As drawn in Figure 5.8, energetic disorder and its accompanying decrease in LUMO and increase in HOMO would make injection of both holes and Anode electrons easier, increasing the number of both carriers. This situation will clearly HOMO vary when the polymer or electrodes are FIGURE 5.8 The lines in this drawing represent energy changed, but the cartoon will work for highlevels of the anode, cathode, HOMO and LUMO levels lighting the important variables the model of a material. The boxes are drawn in to represent the takes into account. While carrier injection is dispersion in HOMO and LUMO levels that can occur facilitated in this cartoon, the distribution of when the conjugation length of a polymer chain or energy levels in the LUMO and HOMO segment varies. In the model studies done by Galvin, mean that some of the carriers are being the solid HOMO and LUMO lines would represent the trapped close to the electrodes. As a result energy levels for the matrix 5PV, while the boxes repreof this trapping, the effective mobility of the sent the energy levels for the PPV added in the blend. carriers and the internal field vary. For discussion purposes, assume that the electron traps are much deeper than the hole traps, electron mobility is lower than hole mobility and that the PLED is electron-limited. In this case, even though more electrons are injected into the device with energetic disorder, the lower effective mobility of electrons with the presence of deep traps could move recombination closer to the cathode, which efficiently quenches luminescence. If the recombination zone is moved enough, PLED efficiency could drop even though more electrons were injected into the device. Using their model, the authors were able to explain the results obtained by Galvin on the effect of blending PPV into 5PV. The model, however, points out that energetic disorder does not always have to hurt device performance. They discuss a device in which hole mobility is significantly higher than electron mobility so that the recombination zone is close to the cathode and much of the luminescence is quenched. If deep hole traps were added to this device, the recombination zone could be moved away from the cathode and the device efficiency would increase. As the field moves forward, this model may prove to be useful for allowing device designers to calculate the type and depth of traps they should add to improve performance. Thus, far we have discussed several variables important to the performance of a polymer used as the active layer of a PLED. The chemical structure of the polymer will determine the PL efficiency and the position of the HOMO and LUMO levels, which are critical to carrier injection. Copolymerization is an alternate scheme for tuning energy levels and, therefore, the color of emitted light. Studies show that the sequence distribution of the monomer units, random, alternating or block, are important for the performance of the polymer in a device. A distribution in chain lengths can result in distribution in conjugation lengths and energy levels, seriously affecting device performance. The packing of polymer units, which is influenced by the solvent used for spinning, affects PL yield and carrier transport. The importance of these variables can also be seen in PV oligomers. A quick review of oligomer studies highlights how sensitive these systems are. LUMO
Cathode
5.2
One-Dimensional Small PPV-Based Molecular Model Compounds for the Study of Polymers
Small linear molecules have been used as models for the study of their corresponding polymer analogs [82–90]. Unlike polymers, small molecules are well defined, exhibit high purity, and thus allow for a more precise study of structure–property relationships. Linear PPV and PPV-derivatized molecules have been a focus of this study to gain further insight into the structural and luminescent properties of
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PPV-based polymers. With the addition of substituents and their placement along the backbone, dramatic changes in optical, thermal, and packing properties are observed. Several studies on small molecules used for this purpose have been reported in recent years and are reviewed below.
5.2.1 Effect of Substituent Addition: Changes in Morphology [85–88] Hadziioannou and coworkers have reported on the synthesis and characterization of an unsubstituted PPV pentamer (1 in Figure 5.9) [85], and dioctyloxy-substituted [86,88], and cyanosubstituted (3 in Figure 5.9) PPV derivatives [87,88]. These provide a good example of how the study of a group of small molecules can provide information on polymer properties. Compounds 1, 2, and 3 are synthesized via the Heck, Wittig, and Knoevenagel condensations, respectively. In this group of compounds, significant changes in packing are observed with the addition of substituents. Single crystals of 1 pack in a herringbone-type arrangement with the structure proposed for unsubstituted PPV [85]. However, a study of single-crystal PL on 2 and 3 reveal changes in packing. Molecule 2 shows vibronic features reminiscent of alkyloxy-substituted PPV. These features are absent in the spectra of the more closely packed compound 3, which is red-shifted, broad, and featureless, characteristic of cyano-substituted PPV [86]. Additionally, for 3, single-crystal studies reveal an interplanar distance of 3.5 A˚, indicating strong p–p interaction between adjacent molecules induced by the cyano moieties [87]. Interactions such as these do not bode well for the use of 1 to 3 in LED applications, however, the insight they provide about their polymer analogs demonstrate their benefit.
5.2.2 Effect of Substituent Placement: Changes in Optical Properties [82] Positional placement of substituents can also alter properties [82,83,91]. This is demonstrated in a study of trimers 4 and 5 (Figure 5.9), reported by Hanack and coworkers [82]. In these, the position of the cyano moiety is changed from a position and b position with respect to the center phenyl ring. In their syntheses, a mono- or bis-functionalized cyanomethylene is reacted with the corresponding mono- or bis-functionalized aromatic aldehyde in a Knoevenagel condensation. With the change in the placement of the cyano moieties from the a to b positions, electronic and optical properties are easily tuned. The solution absorption and emission in film red shifts by more than 50 and 100 nm, respectively, as the position of the cyano units changes from the a to b. Small modifications may be implemented in polymer analogs to tune color output.
1 C6H13 O
2
CN
C8H17O
4 OC8H17 C8H17O
3
NC
CN
O
C6H13 O
C6H13 CN
CN CN
FIGURE 5.9
OC8H17
Small molecules used to study structure–property relationships in PPV systems.
O
C6H13
5
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Conjugated Polymers: Processing and Applications
Two-Dimensional Small PPV-Based Molecules: Effects on Charge Delocalization
Efficient and balanced charge transport is essential to the optimization of molecular and polymeric light-emitting diodes [92]. As discussed previously, in oligomeric and polymeric systems, charge transport occurs via two routes. In intrachain transport, charges are delocalized across the polymer’s linear conjugated backbone. As a result, the more planar and defect-free the backbone is, the more extended the conjugation is. In the process of interchain transport, charges move from chain to chain through a hopping mechanism. In some of these oligomer and polymer systems, two-dimensional charge delocalization can be achieved through p–p stacking between adjacent chains. However, this interaction also increases their tendency to crystallize, leading to the formation of excimers, which quench PL in films. Molecules that have the potential to intrinsically delocalize charge in two dimensions, while avoiding crystallization are, therefore, excellent candidates for improved emissive materials in PLEDs. Liquid crystal and particularly discotic liquid crystal materials present the opportunity to make molecules with p–p stacking that extends in two or three dimensions [93–95]. Liquid crystals can be engineered to be solution-processable, but adopt phases in which the morphology promotes significant molecular overlap and p–p stacking, known to increase carrier mobility [94,96,97]. Thus, this approach has mostly been used in well-defined small molecule systems; however examples of this approach appear in polymeric systems as well [98].
5.3.1
Intrinsic Two-Dimensional Charge Delocalization in PPV Derivative
The polymer shown in Figure 5.10 is comprised of a ladder-type tri(p-phenylene) (LPP) and oligo (phenylenevinylene) (OPV) segments with an orthogonal structure that allows for some degree of intrinsic two-dimensional charge delocalization between those two segments [98]. In solution, as expected, the absorption spectrum of 6 (Figure 5.10) shows features of both the LPP and OPV portions of the polymer, however, the emission spectrum shows only emission from the OPV fragments, indicating a charge transfer mechanism from the LPP to OPV segments. Single-layer devices, having the ITO=6=Ca=Ag configuration, showed blue light emission (m ¼ 1), with a turn-on voltage of 12 V, brightness of 6.6 cd=m2 at current density of 1 mA=mm2, and efficiency of 0.005%. The poor performance of the device is attributed to inefficient electron injection, which can be improved with the incorporation of hole-transporting layers into the device architecture. Though not designed for the purpose of intrinsic two-dimensional charge delocalization, the energy transfer mechanism evident in the polymer provides some indication that this process may be present. Compounds 10 and 11, shown in Figure 5.11, with four conjugated arms attached to a nonconjugated central sp3 hybridized carbon core have recently been reported as emissive materials for LEDs by Bazan and colleagues [99,100]. The syntheses of these molecules, also shown in Figure 5.11, follow
C12H25
6 ]
n
[[ ]m
C12H25
FIGURE 5.10
Ladder polymer that permits some delocalization in two dimensions.
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9 (H3C)3C
C
C (CH3)3
8
C(CH3)3
7 10
11
(H3C)3C
FIGURE 5.11 [99, 100].
C(CH3)3
Four-arm phenylenevinylene star with a nonconjugated core as prepared by Bazan and colleagues
typical Heck reaction conditions. Molecule 10 is prepared in high yield (86%) via the palladiumcatalyzed coupling of tetrakis(4-iodophenyl)methane (7) with the vinyl-terminated styrene (8). Similarly, compound 11, with its longer arms, is made from the coupling of 7 with 4,40 -tert-butylvinylstilbene (9) to yield the pale yellow product (percent yield 31%). In these low molecular weight, well-defined, pure molecules, structure plays a very important role in morphology and subsequently optical properties. For 11, the tetrahedral geometry with the long conjugated arm minimizes p–p interactions and favors the formation of solid amorphous powders and films. In contrast, 10, with shorter arms, is obtained as single crystals from a methanol–benzene mixture. Differential scanning calorimetry (DSC) analysis also points to the crystalline nature of 10, showing a sharp melting transition at 2748C. Molecule 11 reveals no such change; rather a step attributed to a glass transition is seen in the thermogram at 1758C. The absorption and emission maxima for compound 10 are seen at 322 and 379 nm, compared to the redshifted peaks of 11 at 368 and 427 nm. Furthermore, both 10 and 11 exhibit red-shifted emissions compared to linear analogs of their arms, indicting some charge delocalization through the tetrahedral core. By allowing control over morphology and increasing the conjugation length of the arms, these materials show potential for use in optoelectronic applications and improvements in charge delocalization.
5.3.2 PPV-Based Dendrimers Dendrimeric PPV-based molecules have been recently studied as novel emissive and transport layers in LEDs [101–104]. Several synthetic approaches have been utilized including the iterative convergent procedure reported by Burn and colleagues [101,102]. In the iterative convergent synthetic method, the dendrons and core are synthesized separately and then coupled. The preparation of the dendrons involves two convergent steps, the Heck followed by Wittig reactions, resulting in dendrons functionalized by an aldehyde. The Heck followed by Stille coupling reactions can also be used but these coupling reactions result in a lower yield of a bromine-functionalized dendron product. Next, the distyrylbenzene core dendrimer (12), shown in Figure 5.12, is synthesized in one of two ways [105]. The first way involves the coupling of p-divinylbenzene with the bromine-terminated dendron again in a Heck reaction. Alternately, the aldehyde-terminated dendron can be reacted under the Wadsworth–Emmons conditions with a bis-phosphonate to again yield the final dendrimer product. In addition to their simple syntheses, another advantage afforded by these molecules is that they allow for color tuning with variations in the structure of the core molecule. The distyrylbenzene core
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t-Bu
t-Bu t-Bu
t-Bu
t-Bu
t-Bu
t-Bu
t-Bu t-Bu
t-Bu
t-Bu
t-Bu
t-Bu t-Bu
FIGURE 5.12
t-Bu 12
t-Bu
PPV-based dendimer with a conjugated core that acts as a trap and the primary emissive species.
highlighted here emits in the blue (440 nm), while changing the core to distyrylanthracene or a porphyry can change the emission to green (580 nm) or red (660 and 720nm), respectively [102]. The reason for this is that the stablemen dendrons, with solubilizing tertiary butyl groups, serve as transporting moieties, which transfer charge to the core. Due to the meta-linkages, which contribute to a break in conjugation, the stilbene units are independent of the fluorescent core. The core acts as a trap and the primary emissive species. Therefore, changing the core’s structure can tune the electronic properties. Single-layered LEDs, with an ITO=12=Ca architecture, performed reasonably well, exhibiting a device efficiency of 0.09%. It should be noted that with this approach, because of breaks in the intrinsic charge delocalization, the electronic and processing properties can be controlled independently. This as well the respectable device efficiency renders this class of molecules as attractive novel candidates for LED applications.
5.3.3
Two-Dimensional Conjugated Molecules with Four Arms and a Conjugated Core
The work described in this section, reported by our group, differs from those already discussed in that for molecules 13–16 conjugation is maintained along the arms as well as through the para- and, to a lesser degree, ortho-linkages at the core [106]. Study of the optical properties of smaller model compounds determined that the tetra-substituted molecule exhibited a high solution quantum yield and potential for improved charge carrier mobility through two- and possibly three-dimensions compared to analogous meta-substituted compounds [107]. The solution-processable compounds 13–16 point to the synthetic versatility of this class of molecules, highlighted by the incorporation of oxadiazole and cyano moieties that demonstrate color tuning of emission. This flexibility enables the chemist to engineer the HOMO and LUMO levels of the molecules. For 13–16, which emit in the yellow green, an orthogonal approach [108,109] employing the Heck and Horner–Emmon reactions alternately was used. The three-part synthesis consists of a dibromo–bisphosphonate core, a vinyl-terminated arm and an aldehyde-terminated arm. For molecule 14, an orange light emitter, a slightly different method is used that involves the coupling of a core molecule to a cyano-terminated arm using the Knoevenagel condensation.
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As a whole, compounds 13–16 show a red-shifted film emission compared to a linear analog, Stokes shift in the 0.7 eV range, and respectable solution PL quantum yield of 0.4–0.8. Significant p–p interactions are observed for 13–16 in concentrated solution and powder samples explaining, in part, the red-shifted emission, though fairly crystalline free films are able to be formed with spin-casting. Even with a lack of p–p stacking, as in 13, the red shift is observed, indicating that in these molecules two-dimensional and possibly three-dimensional charge delocalizations may be present. The existing of two-dimensional delocalization is supported by studies on these molecules performed by Seibbeles [110]. Multilayer LEDs composed with a ITO=PEDOT:PSS=Emissive two-dimensional molecule=LiF=Al device architecture showed promising efficiencies of 0.1% and 0.2% for 14 and 16, respectively, revealing the promising use of this class of compounds for emissive applications and improved charge delocalization (Figure 5.13). Recently, Wilson et al. have reported on the synthesis and characterization of poly(p-phenyleneethynylene) (PPE) and PPV hybrid oligomers arranged orthogonally to each other [111]. Using the Horner– Sonogashira reactions, the two-dimensional molecules have been synthesized with varying R groups, two of which are highlighted here (17 and 18 in Figure 5.14). Optical properties are greatly varied depending on chemistry of side groups. For molecule 17, high-solution fluorescent yields of 78% and 88% are observed in both hexanes and chloroform, respectively. However, for 18, fluorescence is dramatically reduced going from 53% in hexanes to only 14% in chloroform. Ground and excited state charge delocalizations are also altered by the nature of the side groups. For 18, the HOMO is calculated to reside only on the PPV arms, whereas the LUMO is located on the PPE segments. Conversely, in molecule 17, this separation between orbitals is not seen in the excited state, H3C
H17C8O
OC8H17
H17C8O H3C
OC8H17
H17C8O
OC8H17
H17C8O
OC8H17 OC8H17
13
H17C8O
CH3
CN
NC
CN
NC
H17C8O H3C
H3C
OC8H17
CH3
H17C8O
H3C
14 OC8H17
N N O
N N O
H17C8O CH3
OC8H17 H17C8O
H3C
N N O
CH3
H3C CH3
OC8H17
H17C8O
OC8H17
H17C8O
16
15 H3C
OC8H17
O N N
H3C
OC8H17
H17C8O CH3
CH3
FIGURE 5.13 Four-arm phenylenevinylene-type based that can be used as an active layer in LEDs. (From Niazimbetova, Z.I., et al., J. phys. chem. B., 108, 8673, 2004 and Niazembetova, Z.I., et al., J. Electromole. chem., 529, 43, 2002.)
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OCH3 N
H3CO
OCH3 F3C
CF3
F3C
CF3
17
18
H3CO N
FIGURE 5.14
Four-arm ethylene stars that can be used for color tuning.
rather the charge is delocalized across the entire molecule. With the versatile nature of the synthesis of these small PPE–PPV hybrids and their potential for color tuning, they present themselves as interesting materials for electronic applications.
5.4
Polyfluorenes
While PPV was the first polymer used in PLEDs and its copolymers and derivatives are still extensively studied, poly(9,90 -dialkyfluorenes) (PFs) have emerged as a prominent commercially viable class. These polymers were of interest early on in the development of PLEDs because the rigidity of the biphenyl group, and the ensuing twisting along the chain, increased the band gap of the polymer, making them emit blue light, which was difficult to get from PPVs. Another advantage of PFs over PPVs is that they are less susceptible to oxidation than PPVs, in which carbonyl formation at the vinyl bond can be a serious problem, as discussed previously. The ability to substitute the methylene bridge with a wide variety of groups allows the chemist to impart solubility, tune the color of emission, or even impart liquid crystallinity [112,113] or chirality [114] to the polymer. PFs are generally thermally and chemically stable, although we will see later that reaction at the methylene bridge can be problematic. This section will discuss some of the important properties of PFs as they relate to their use in PLEDs. It will not, however, begin to review the field of PF synthesis, characterization, or utilization in PLEDs. To obtain a more comprehensive picture, the reader is referred to other reviews of PFs [6,11,114–117]. Not long after the first report of PPV-based PLEDs, Ohmori et al. [118] reported on the first PLED using a PF active layer, which emitted blue light. PFs can be synthesized via a variety of techniques including Grignard coupling, Heck coupling, Stille coupling, Yamamoto coupling [119], and Suzuki coupling [117]. Scientists at Dow Chemical have developed a wide variety of fluorene homo- and copolymers that emit in the blue, green, and red regions of the spectrum. Using a modified Suzuki coupling (Figure 5.15), these researchers have been able to get high molecular weight (500,000 Da) polymers with reasonable polydispersities [6]. This approach enables the synthesis of a plethora of copolymers by substituting the dibromofluorene with an alternate aromatic dibromide. Some of the alternate aromatic units used in copolymerization include carbazole [6,120], thiophenes [121], bithiophenes [121], oxadiazoles [122], quinoline [123], quinoxaline [123], phenylenecyanovinylene [123], triarylpyrazoline [124], anthracenes [6], phenyls [6], aromatic amines [6,125], benzothiadiazole [126–128], and binapthyl [129]. The inclusion of co-monomers can do more than just tuning the color
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O
O B O
Br
+ Br
B O R
R
R
R
Pd0 Na2CO3 Toluene and H2O *
*
R R Poly(9,9-diaklyfluorene) Br
Br NN
Triarylpyrazoline
N Br
S
Br
N Br
Benzothiadiazole
Br N
Triaryl amine
FIGURE 5.15 Synthetic scheme for the synthesis of polyfluorenes. Note that the dibromofluorene can be fully or partially replaced with other aromatic dibromides, like the ones drawn here, to synthesize copolymers containing fluorenes.
of the emitted light. Incorporation of aromatic amines has been shown to improve hole injection by lowering the oxidation potential of the polymer [130]. Copolymers with aromatic amines had TOF mobilities that were an order of magnitude higher than those of the parent homopolymer [131,132]. Cao and colleagues have been able to form efficient electron injection layers from a copolymer of fluorene–benzothiadiazole in which some of the alkyl chains on the fluorenes have quarternized ammonium end groups [133]. Dow has done considerable research on color tuning and studying the lifetimes of PF-based PLEDs. In a recent review [6], they note that in a ITO=PEDOT:PSS=light-emitting polymer=Ca=Ag or Al device, green PLEDs have an efficiency of 7.6 lm=W. The red devices can obtain 1.4 lm=W and the blue 1.3 lm=W, all at 200 Cd=m2. The lifetime of the blue device was, however, only 1500 h. The lifetime of blue-emitting diodes is one of the major hurdles faced in the commercialization of full color PLED displays, which require tens of thousands of hours. In general, the PFs have good thermal stability with glass transitions in the range of 1108C [6]. PFs with 2-ethylhexyl alkyl groups or dioctyl groups on the methylene carbon can form liquid crystalline phases and have an energy gap of 3 eV. If films of these PFs are spun from solution and then thermally annealed, a new absorption peak around 2.8 eV appears [134]. This peak is attributed to increased conjugation that results from planarization of the chains in a new b phase of the polymer. PFs can emit a beautiful blue color enabling the fabrication of excellent full color displays, but their major problem is instability in the color of the emitted light. During operation of a PLED or during thermal annealing of the PF film, green light begins to be emitted, destroying the color purity. Originally it was believed that the appearance of green emission, centered at 530 nm, was the result of excimer [135–137] or aggregate [138] formation. Chemists, therefore, tried to eliminate the appearance of green emission by preventing eximer or aggregate formation through attaching dendritic [139,140] side chains or end-capping the chains [137]. These approaches improved the color stability, but did not completely eliminate the problem. Recently researchers have begun to think that the origin of the green emission is due to the presence of emissive keto units [115,141,142] (Figure 5.16) in the chain. List has conducted a beautiful set of experiments reviewed in a recent publication [134] that lend credence to this theory. List notes that when a PF film is heated or exposed to UV light in the presence of oxygen the peaks at 2.96 and 2.78 eV, which correspond to the blue-emitting species, decrease, whereas there is a
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*
* R
R
y
x O
n
FIGURE 5.16 Polyfluorene with fluorenone units. The presence of these keto groups is now believed to be the origin of green emission from PF-based PLEDs. The energy level of the fluorenone is lower than that of the fluorene, which, therefore, serves to funnel energy down the chain to the green-emitting keto.
concurrent increase in the growth of the green peak at 2.3 eV. Using IR, List and coworkers were able to correlate these spectral changes to the growth of a new carbonyl stretch in the IR, attributed to the presence of keto groups. These same changes occurred in a film when it was used as the active layer of a device. Copolymerization of 9,9-diakylfluorene monomers with the keto-containing fluorenone monomer supported the conclusion that the green emission is due to the presence of keto groups [134]. Elimination of the keto groups is difficult. They can form during device operation, especially if oxygen is present, but some of the keto units were found to have been introduced into the polymer during synthesis. Mono-substituted fluorenes are very susceptible to oxidation to the keto unit during polymerization. Rigorous exclusion of these impurities during synthesis does improve PF performance [143], but is not a perfect solution. To eliminate the formation of fluorenone units, two groups have independently developed new polymers in which the methylene carbon is replaced [7]. A group from Durham replaced the methylene with SO2 groups [7]. Holmes and his colleagues prepared 2,7-dibenzosilole and synthesized homo- and copolymers of the silole monomer with dialkyl fluorene and phenyl end-caps [144]. The silole homopolymer had a band gap of 2.93 eV and maintained its color emission even when heated to 2508C for 16 h. These approaches indicate that PFs may one day form the basis for a polymeric full color display.
5.5
Phosphorescence in PLEDs
Most of the PLEDs fabricated to date have relied on fluorescence, the decay of singlet states, for light generation. In small molecule light-emitting diodes (SMOLEDs), it has been well established that the spin statistics limit electroluminescence to at best 25% of the PL of a film. This limit is determined by the fact that there are three triplet states and one singlet state that can be formed when the hole and the electron meet, since these injected carriers have uncorrelated spins. As Thompson and Forrest have pointed out this means that three-fourths of the potential light is being wasted since triplets generally decay nonradiatively. To overcome this limitation, they have focused considerable effort on fabricating devices (PHOLEDs) [145–147] based on phosphorescence from organometallic complexes, often based on iridium or platinum, in which spin–orbit coupling allows efficient radiative decay of the triplet. In some of these devices, the emission can come from both the singlet and triplet states, significantly increasing the efficiency. Whether these spin statistics apply to polymeric devices is still a matter of considerable debate. Baldo and colleagues have done measurements on singlet–triplet ratios in MEH– PPV, which indicate that the 3:1 triplet to singlet ratio holds [148]. Bredas has, however, done calculations, which show as the chain length increases the 25% EL to PL rule breaks down [149–151]. These calculations support experimental work in which EL to PL ratios higher than 25% have been found [152–154]. For a detailed discussion of their topic the reader is referred to an excellent review of fluorescence and phosphorescence in organic materials written by Kohler et al. [153]. While the debate about whether the 25% EL to PL rule applies to polymers continues, some researchers have begun to fabricate electrophosphorescent polymer devices (PPHOLEDs). A comprehensive review of these devices is beyond the scope of this chapter, but it is worth mentioning a few of the different approaches being taken by researchers. In the simplest format, phosphorescent complexes
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Polymers for Use in Polymeric Light-Emitting Diodes t - Bu
O Ir
Ir O
N
N
3
2
Ir(Bu-ppy)3
FIGURE 5.17
(ppy)2Ir acac
Structures of iridium complexes used in some electrophosphorescent devices.
often based on iridium, as seen in Figure 5.17, are blended with a polymer capable of transporting holes and electrons. To obtain reasonable efficiencies, the lowest lying triplet state of the polymer should match the energy levels of the phosphor and aggregation of the phosphor must be avoided. In some cases, the polymer used for blending is a conjugated polymer [155–157] as in the research done by Cao et al. [158], where soluble poly(p-phenylenes) (PPP) were used as the matrix. The best external quantum efficiency obtained by Cao was 3% when 2% of Ir(Bu-ppy)3 was incorporated into a cyano-substituted PPP. If the iridium complex was not used, the device had an efficiency of only 0.22%, highlighting the advantages of using phosphorescence. Wu et al. [125] used fluorene copolymers that either had hole-transporting triaryl amine moieties in the main chain and electron-transporting oxadiazoles as pendant groups on the fluorene units or alternatively had a fluorene with oxadiazole side chains copolymerized with a fluorene with pendant aryl amines. When blended with an osmium phosphor, the PPHOLED had an external quantum efficiency of 9.3%. An alternate approach has been to use a nonconjugated polymer as the matrix [159]. Most typically, poly(vinyl carbazole) (PVK), a hole-transporting polymer, is blended with an electron-transporting compound, often 2-tert-butylphenyl-5-biphenyl-1,3,4-oxadiazole (PBD) (Figure 5.18). A phosphor is then added to a solution of PVK and PBD and a film is spun as the active layer of a PPHOLED. While this approach can result in mediocre performance, Yang using commercially available iridium complexes has obtained devices with power efficiencies of 3 lm=W [160], Gong, using a fluorene-based iridium, obtained external quantum efficiencies of 8% [161] and Chang with another iridium complex got red emission with a luminance efficiency of 8.8 cd=A [162]. Frechet [163] used an alternate matrix to PVK:PBD. He copolymerized vinyl oxadiazoles with vinyl amines and used this as the matrix. When
* *
n N
O N N
PVK
PBD
FIGURE 5.18 PVK is a hole-transporting polymer commonly mixed with the electron transport molecule PBD and used as a host in electrophosphorescent devices.
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N Ir O
O F3C C8H17
C10H20
C8H17
N
C8H17
C8H17 *
*
*
y X 19 Br
C8H17
C8H17 S N Ir
40
N S C8H17 20
FIGURE 5.19
Br C8H17 40
Polymers incorporating phosphors into their chains in order to prevent aggregation.
mixed with a platinum phosphor and fabricated into a PPHOLED, the device had 4.6% external quantum efficiency. A problem with blending the phosphor into a matrix is that phase separation can lead to aggregation of the phosphor, which results in low phosphorescent yields. To avoid this problem, a few researchers have incorporated iridium complexes into a polymer. A few of these structures are shown in Figure 5.19. When Jiang et al. [120] mixed polymer 19 with PBD, the PPHOLEDs had an external quantum efficiency of 3.4% and a luminance efficiency of 2.9 cd=A. Holmes [164] made a series of fluorene polymers that incorporated iridium phosphors. The most efficient PPHOLEDs had an external quantum efficiency of 1.5%, which came from triplet emission and was fabricated with 20 as the active layer. These polymers prove that incorporation of the phosphors into a polymer chain can produce efficient devices and will likely be an area of further research. The CDT Web site notes that dendimers can also be used for this purpose. Dendrons can prevent the phosphorescent core from aggregating and thus, reducing PL [165]. There is little doubt that research in the arena of electrophosphorescent displays will increase in the future.
5.6
Summary
Significant progress in the development of electroactive polymers for use as the active layer in PLEDs has been made since the discovery by Friend and Holmes that PPV could be used in a light-emitting diode. Some commercial products now use PLEDs for their displays and researchers and corporations are
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considering using PLEDs for solid state lighting applications. Solution processing and color tuning have become commonplace. The classes of polymers being considered for use has expanded well beyond PPVs, with polyfluorenes showing great potential especially when blue light is needed. The number of homopolymers and copolymers synthesized and evaluated is so large, that a comprehensive review of all these materials would fill a book. Despite all this progress, however, there remains a considerable amount of research to be done. Scientists are just beginning to understand the fundamental structure–property relations that govern the performance of this class of materials. As seen with the fluorenone groups in PFs and the carbonyls in PPVs, trace amounts of impurities can have profound impact on the performance of PLEDs. The extent to which energetic disorder, or variations in conjugation length, matter is just beginning to be explored. The balance that must be struck between chain separation increasing PL, but decreasing mobility and perhaps the number of carriers is only understood qualitatively. While researchers have begun to look at delocalizing charge in two dimensions, they are just scratching the surface of this topic. More detailed studies of the impact of copolymer sequence distribution on device performance are needed before polymers can be designed from first principles. Before printable devices become commonplace, researchers will need to make considerable progress on injecting electrons from environmentally stable cathodes and a significant amount of physics remains to be understood before PPHOLEDs become commercialized. In summary, the field still holds many challenges and opportunities for researchers.
References 1. Chiang, C.K., C.R. Fincher, Y.W. Park, A.J. Heeger, H. Shirakawa, E.J. Louis, S.C. Gau, and A.G. Macdiarmid.1997. Electrical-conductivity in doped polyacetylene. Phys Rev Lett 39 (17): 1098–1101. 2. Shirakawa, H., E.J. Louis, A.G. Macdiarmid,C.K. Chiang, and A.J. Heeger. 1997. Synthesis of electrically conducting organic polymers–halogen derivatives of polyacetylene. (Ch)X. J Chem Soc-Chem Commun 578–580. 3. Burroughes, J., D. Bradlet, A. Brown, R. Marks, K. Mackay, P. Burn, and A. Holmes. 1990. Lightemitting diodes based on conjugated polymers. Nature 347:539–541. 4. Phillips, polymer light-emitting diodes. http:==Www.Research.Philips.Com=Technologies=Display= Polyled=Polyled= 5. Shimoda, T., K. Morii, S. Seki, and H. Kiguchi. 2003. Inkjet printing of light emitting polymer displays. Mrs Bulletin 821–827. 6. Wu, W.S., M. Inbasekaran, M. Hudack, D. Welsh, W.L. Yu, Y. Cheng, C. Wang, S. Kram, M. Tacey, M. Bernius, R. Fletcher, K. Kiszka, S. Munger, and J. O’brien. 2004. Recent development of polyfluorene-based Rgb materials for light emitting diodes. Microelectron J 35: (4): 343–348. 7. Milgrom, L. 2005. Blue light shines on polymer leds. Chem World 2 (7): 14. 8. Grem, G., and G. Leising. 1993. Electroluminescence of wide-bandgap chemically tunable cyclic conjugated polymers. Synth Met 57 (1): 4105–4110. 9. Grem, G., G. Leditzky, B. Ullrich, and G. Leising. 1992. Blue electroluminescent device based on a conjugated polymer. Synth Met 51 (1–3): 383–389. 10. Grem, G., G. Leditzky, B. Ullrich, and G. Leising. 1992. Realization of a blue-light-emitting device using poly(para-phenylene). Adv Mater 4 (1): 36–37. 11. Bernius, M., M. Inbasekaran, E. Woo, W.S. Wu, and L. Wujkowski. 2000. Fluorene-based polymerspreparation and applications. J Mater Sci Mater Electron 11 (2): 111–116. 12. Grice, A.W., D.D.C. Bradley, M.T. Bernius, M. Inbasekaran, W.W. Wu, and E.P. Woo. 1999. High brightness and efficiency blue light-emitting polymer diodes. Appl Phys Lett 73 (5): 629–631. 13. Roncali, J. 1992. Conjugated poly(thiophenes)—synthesis, functionalization, and applications. Chem Rev 92 (4): 711–738.
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135. Lee, J.I., G. Klaerner, and R.D. Miller. 1999. Structure-property relationship for excimer formation in poly(alkylfluorene) derivatives. Synth Met 101 (1–3): 126–126. 136. Bliznyuk, V.N., S.A. Carter, J.C. Scott, G. Klarner, R.D. Miller, and D.C. Miller. 1999. Electrical and photoinduced degradation of polyfluorene based films and light-emitting devices. Macromolecules 32 (2): 361–369. 137. Lee, J.I., G. Klaerner, and R.D. Miller. 1999. Oxidative stability and its effect on the photoluminescence of poly(fluorene) derivatives: End group effects. Chem Mater 11 (4): 1083–1088. 138. Lemmer, U., S. Heun, R.F. Mahrt, U. Scherf, M. Hopmeier, U. Siegner, E.O. Gobel, K. Mullen, and H. Bassler. 1995. Aggregate fluorescence in conjugated polymers. Chem Phys Lett 240 (4): 373–378. 139. Pogantsch, A., F.P. Wenzl, E.J.W. List, G. Leising, A.C. Grimsdale, and K. Mullen. 2002. Polyfluorenes with dendron side chains as the active materials for polymer light-emitting devices. Adv Mater 14 (15): 1061–1064. 140. Setayesh, S., A.C. Grimsdale, T. Weil, V. Enkelmann, K. Mullen, F. Meghdadi, E.J.W. List, and G. Leising. 2001. Polyfluorenes with polyphenylene dendron side chains: Toward non-aggregating, light-emitting polymers. J Am Chem Soc 123 (5): 946–953. 141. Lupton, J.M., M.R. Craig, and E.W. Meijer. 2002. On-chain defect emission in electroluminescent polyfluorenes. Appl Phys Lett 80 (24): 4489–4491. 142. List, E.J.W., R. Guentner, P.S. De Freitas, and U. Scherf. 2002. The effect of keto defect sites on the emission properties of polyfluorene-type materials. Adv Mater 14 (5): 374–378. 143. Craig, M.R., M.M. De Kok, J.W. Hofstraat, A. Schenning, and E.W. Meijer. 2003. Improving color purity and stability in a blue emitting polyfluorene by monomer purification. J Mater Chem 13 (12): 2861–2862. 144. Chan, K.L., M.J. Mckiernan, C.R. Towns, and A.B. Holmes. 2005. Poly(2,7-dibenzosilole): A blue light emitting polymer. J Am Chem Soc 127 (21): 7662–7663. 145. Adamovich, V., J. Brooks, A. Tamayo., A.M. Alexander, P.I. Djurovich, B.W. D’andrade, C. Adachi, S.R. Forrest, and M.E. Thompson. 2002. High efficiency single dopant white electrophosphorescent light emitting diodes. New J Chem 26 (9): 1171–1178. 146. Kohle, A., J.S. Wilson, and R.H. Friend. 2002. Fluorescence and phosphorescence in organic materials. Adv Mater 14 (10):701–707. 147. Adachi, C., M.A. Baldo, M.E. Thompson, and S.R. Forrest. 2001. Nearly 100% internal phosphorescence efficiency in an organic light-emitting device. J Appl Phys 90 (10): 5048–5051. 148. Segal, M., M.A. Baldo, R.J. Holmes, S.R. Forrest, and Z.G. Soos. 2003. Excitonic singlet–triplet ratios in molecular and polymeric organic materials. Phys Rev B 68:(7). 149. Shuai, Z., D. Beljonne, R.J. Silbey, and J.L. Bredas. 2000. Singlet and triplet exciton formation rates in conjugated polymer light-emitting diodes. Phys Rev Lett 84 (1): 131–134. 150. Beljonne, D., A.J. Ye, Z. Shuai, and J.L. Bredas. 2004. Chain-length dependence of singlet and triplet exciton formation rates in organic light-emitting diodes. Adv Funct Mater 14 (7): 684–692. 151. Beljonne, D., Z.G. Shuai, A.J. Ye, and J.L. Bredas. 2005. Charge-recombination processes in oligomer-and polymer-based light-emitting diodes: A molecular picture. J Soc For Information Display 13 (5): 419–427. 152. Cao, Y., I.D. Parker, G. Yu, C. Zhang, and A.J. Heeger. 1999. Improved quantum efficiency for electroluminescence in semiconducting polymers. Nature 397 (6718): 414–417. 153. Kohler, A., J.S. Wilson, and R.H. Friend. 2002. Fluorescence and phosphorescence in organic materials. Adv Eng Mater 4 (7): 453–459. 154. Wohlgenannt, M., and Z.V. Vardeny. 2003. Spin-dependent exciton formation rates in Pi-conjugated materials. J Phys Condensed Matter 15 (3): R83–R107. 155. Campbell, L.H., D.L. Smith, S. Tretiak, R.L. Martin, C.J. Neef, and J.R. Ferraris. 2002. Excitation transfer processes in a phosphor-doped poly(p-phenylene vinylene) light-emitting diode. Phys Rev B 65:(8).
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156. Jiang, C.Y., W. Yang, J.B. Peng, S. Xiao, and Y. Cao. 2004. High-efficiency, saturated red-phosphorescent polymer light-emitting diodes based on conjugated and non-conjugated polymers doped with an Ir complex. Adv Mater 16 (6): 537–541. 157. Gong, X., W.L. Ma, J.C. Ostrowski, K. Bechgaard, G.C. Bazan, A.J. Heeger, S. Xiao, and D. Moses. 2004. End-capping as a method for improving carrier injection in electrophosphorescent lightemitting diodes. Adv Funct Mater 14 (4): 393–397. 158. Zhu, W., Y. Mo, M. Yuan, W. Yang, and Y. Cao. 2002. Highly efficient electrophosphorescent devices based on conjugated polymers doped with iridium complexes. Appl Phys Lett 80 (12): 2045–2047. 159. Chen, F.C., S.C. Chang, G.F. He, S. Pyo, Y. Yang, M. Kurotaki, and J. Kido. 2003. Energy transfer and triplet exciton confinement in polymeric electrophosphorescent devices. J Polym Sc Part B Polym Phys 41 (21): 2681–2690. 160. Yang, X.H., D. Neher, D. Hertel, and T.K. Daubler. 2004. Highly efficient single-layer polymer electrophosphorescent devices. Adv Mater 16 (2): 161–166. 161. Gong, X., J.C. Ostrowski, D. Moses, G. C Bazan, and A.J. Heeger. 2003. Electrophosphorescence from a polymer guest–host system with an iridium complex as guest: Forster energy transfer and charge trapping. Adv Funct Mater 13 (6): 439–444. 162. Chang, C.Y., S.N. Hsieh, T.C. Wen, T.F. Guo, and C.H. Cheng. 2006. High efficiency red electrophosphorescent polymer light-emitting diode. Chem Phys Lett 418 (1–3): 50–53. 163. Frechet, J.M.J. 2005. Functional polymers: From plastic electronics to polymer-assisted therapeutics. Prog Polym Sc 30 (8–9): 844–857. 164. Sandee, A.J., C.K. Williams, N.R. Evans, J.E. Davies, C.E. Boothby, A. Kohler, R.H. Friend, and A.B. Holmes. 2004. Solution-processable conjugated electrophosphorescent polymers. J Am Chem Soc 126 (22): 7041–7048. 165. Cambridge Display Technology, http:==www.cdtltd.co.uk=technology=40.asp
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6 Organic Electro-Optic Materials 6.1 6.2 6.3 6.4 6.5
Introduction to Electro-Optic Activity ........................... 6-1 Theoretically Inspired Optimization of Molecular First Hyperpolarizability ................................................... 6-5 Characterization of Molecular First Hyperpolarizability............................................................ 6-9 Synthesis of Chromophores ............................................. 6-9 Theoretically Inspired Optimization of Macroscopic Electro-Optic Activity ..................................................... 6-11 Chromophore–Polymer Composite Materials . Dendrimers and Super=Supramolecular Materials=Organic Glasses
6.6 6.7 6.8 6.9 6.10 6.11 6.12 6.13 6.14 6.15
Larry R. Dalton
6.1
6.17
Characterization of Electro-Optic Activity ................... 6-16 Auxiliary Properties and Their Characterization ......... 6-17 Lattice Hardening............................................................ 6-24 Fabrication of Devices and Circuits: Electron Beam, Reactive Ion, and Photo-Etching ....................... 6-25 Fabrication of Devices and Circuits: Soft and Nano-Imprint Lithography............................................. 6-26 Stripline, Ring Microresonator, Spatial Light Modulator, and Photonic Band Gap Devices ............... 6-26 Conformal and Flexible Devices .................................... 6-28 Integration with Silicon Photonics ................................ 6-28 Integration with VLSI Electronics.................................. 6-29 Terahertz Signal Generation and Detection, Optical Rectification, and Electro-Optic Sensors ......... 6-29 Applications and Commercialization ............................ 6-30
Introduction to Electro-Optic Activity
Electro-optic (EO) activity [1–10] typically involves the application of a low-frequency (0–200 GHz) electric field to a material, which results in charge displacement in the material. This charge perturbation, in turn, can alter the velocity of light passing through the material. Another way of thinking about electro-optic materials is to view them as materials whose index of refraction can be varied by application of an applied electric field (voltage). A variety of charge displacement phenomena can give rise to electro-optic activity. These include cooperative molecular reorientation as in liquid crystalline materials where displacement of both electronic and nuclear charges is involved; ion displacement as in ionic crystalline materials such as lithium niobate; and electronic displacement in p-electron molecules 6-1
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appropriately ordered in the solid state. In each of these cases, the effective mass displaced is quite different and larger mass displacement translates into slower response time (and bandwidth limitation for devices fabricated from the respective materials). Another factor that can influence the response of electro-optic materials is disparate velocities of optical and electrical waves copropagating in the material. Lithium niobate has a high dielectric constant and a relatively low index of refraction, leading to significant mismatch in the velocity of propagating electric and optical waves. This limits bandwidth to about 10 GHz for a 1 cm simple device structure; devices with higher bandwidths have been achieved by clever device engineering that circumvents this fundamental limitation. The bandwidth of liquid crystalline electro-optic devices is typically less than 1 MHz limited by reorientation time. The intrinsic bandwidth of p-electron ‘‘electronic’’ electro-optic materials is at least several hundreds of gigahertz (GHz) and 3 dB bandwidths of 200 GHz have been reported [11]. Indeed, the bandwidth of devices fabricated from organic electro-optic materials is typically not limited by any property of the organic material but rather by resistive (micro- and millimeter wave) losses in the metal electrodes used to bring the low-frequency field to the electro-optic material. When organic electro-optic materials are used in terahertz signal (difference frequency or optical rectification) generation or detection, operation can be achieved to tens of terahertz [12,13]. Clearly, one of the advantages afforded by organic electro-optic materials is the fast response of delocalized p-electrons to electric field perturbation; the most fundamental response time is the phase relaxation time of the p-electron system, which is typically on the order of tens of femtoseconds. This most fundamental bandwidth limitation for organic materials is the same for both second (electro-optic) and third (all-optical) order nonlinear optical phenomena. The greatest challenge for the device engineer is finding device (e.g., electrode) designs that permit the full bandwidth potential of organic electro-optic materials to be exploited. All-optical signal processing, using third-order materials, avoids this device engineering challenge but third-order optical nonlinearities are much smaller than second-order nonlinearities. These optical nonlinearities are the coefficients of the second (quadratic in electric field) and third (cubic in electric field) terms in the power series expansion of polarization in terms of electric field. As is typical of power series expansions, the magnitudes of the coefficients decrease progressing to higher order terms. Organic electro-optic materials are most commonly encountered as dipolar charge transfer molecules organized into noncentrosymmetric material lattices, although octupolar [14–16] molecules can also exhibit electro-optic effects. Dipolar charge transfer molecules can be viewed as modular materials consisting of an electron-rich (donor) region separated from an electron-deficient (acceptor) region by a conjugated p-electron bridge (see Figure 6.1). In the two-state or Mulliken picture of charge transfer molecules, two limiting forms are envisioned: Neutral and charge-separated (zwitterionic) forms. In real molecules, the ground and first excited states can be envisioned as mixtures of these two limiting forms and the application of an applied electric field can be viewed as changing the mixing of these forms thus changing the polarization and index of refraction of the material. The parameter relevant to electrooptic activity at the molecular level is the molecular first hyperpolarizability (b), which is the coefficient of the second-order (quadratic) term in the power series expansion of molecular polarization in terms of applied electric field. In the simplest two-state model description, b ¼ (mee mgg)Mge=DEge, where mee mgg is difference in dipole moments of the excited and ground states, Mge is the transition matrix between these states, and DEge is the HOMO–LUMO gap (electronic interband transition energy).
Electron donor
Ra
S N
Rb
Rf
Re S
Rc
Rd
NC
NC CN O
F3C Rg
π-Conjugated bridge
FIGURE 6.1
The components of a typical electro-optic chromophore are shown.
Electron acceptor
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Molecular first hyperpolarizability increases significantly with increasing length of the p-electron bridge; but in general, structure–function relationships for b are not simple and optimization of b requires guidance from carefully evaluated quantum mechanical calculations (see Section 6.2). If the ground state of a chromophore is dominated by the neutral form, the molecular first hyperpolarizability will be of positive sign whereas it will be negative if the chromophore ground state is dominated by the zwitterionic contribution. Obviously, solvatochromic shifts will be of opposite signs for these two cases. Individual molecules will contribute to macroscopic electro-optic activity in an additive manner; hence, the linear dependence on chromophore number density, N, in the master equation for electrooptic activity, r33 ¼ 2N b f (v) hcos3 ui=n4 where r33 is the principle element of the electro-optic tensor, f (v) is a local field factor, hcos3 ui is the acentric order parameter, and n is the material index of refraction. As can be seen from the master equation, noncentrosymmetric (acentric) ordering of dipolar charge transfer chromophores is required for nonzero electro-optic activity. Normally, organic electrooptic thin films are prepared by dissolving organic chromophores in a polymer lattice and electric field poling the resultant material near its glass transition temperature, Tg. Alignment can also be achieved by sequential synthesis=self-assembly protocols [17–21]. In the absence of intermolecular electrostatic interactions, a very simple relationship is predicted for the acentric order parameter under the condition of electric field poling, namely, hcos3 ui ¼ mF=5kT where m is the chromophore dipole moment, F is the electric poling field felt by the chromophore (i.e., the poling field corrected for the presence of the host dielectric medium), k is the Boltzmann constant, and T is the poling temperature in kelvin. In this simple independent particle description, electro-optic activity is predicted to increase in a linear manner with chromophore concentration (number density N ) and with chromophore dipole moment. In reality, chromophores with large b values also tend to exhibit large dipole moments (e.g., greater than 10 D) and intermolecular electrostatic interactions cannot be neglected. We shall see in this chapter that careful consideration of intermolecular electrostatic interactions using statistical mechanical methods is critical to the rational optimization of macroscopic electro-optic activity. Indeed, the engineering of ultimate electro-optic materials is becoming a textbook example of applied nanoscience or nanotechnology as individual chromophores are shaped at the nanometer scale and assembled into nanostructured material lattices. Electro-optic activity is a tensor quantity. For electrically poled dipolar organic chromophores, the two nonzero tensor elements are r33 and r13. r33 is a shorthand notation for r333 where ‘‘333’’ indicates that chromophore (electro-optic tensor principal axis), applied electrical (radiofrequency, microwave, or millimeter wave) field, and optical field (optical mode) principal axes all point in the same direction (e.g., defined laboratory ‘‘z’’ or ‘‘3’’ axis direction). For many materials and poling conditions, r33 3(r13), although quite different results are obtained with techniques such as optically assisted (laser) poling [22,23] applied to chromophores such as charge transfer versions of azobenzene and stilbene molecules. In this case, a polarized optical field (inducing trans–cis–trans isomerization) drives the chromophores to assume an Ising-like lattice and the applied electrical poling field acts to break the optically induced centrosymmetric symmetry. Deviation from the r33 3(r13) relationship is also observed at very high poling fields and for cases where strong intermolecular electrostatic interactions influence poling-induced acentric order. Most device configurations will utilize the r33 component but a few will depend on r13 or r33–r13 [24]. It is interesting to consider the relationship of electro-optic activity in organic p-electron materials to other applications (including organic electronics and photovoltaic behavior) dependent on electrical conductivity (electron and hole mobilities). Electro-optic activity depends upon controlled charge displacement while these other applications depend upon controlled charge transport. Both are facilitated by weakly bound p-electrons. For electro-optic activity, it is critical to confine charge displacement (under the action of an applied electric field) to within the chromophore molecule. In the cases of photovoltaic and organic electronic applications, one wants to promote charge separation and transport. If the donor and acceptor ends of molecules, characterized by large b values, come into close proximity (at high number densities) then very high charge mobilities (e.g., comparable to or greater than amorphous silicon [1–10 cm2=V s]) can result. Careful attention must be given to nanoscopic assembly
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of chromophores to prevent charge migration for realization of optimum electro-optic activity. Conductivity can also contribute to bias voltage drift in the operation of electro-optic devices such as Mach Zehnder modulators. On the other hand, nanoscopic engineering can be turned around to realize effective charge transport and exceptional electronic properties. Later in this chapter, we will discuss how organic electro-optic materials can be used for optical rectification [25] and photodetection. Second harmonic generation (SHG) [26,27] and optical rectification [25] are manifestations of second-order optical nonlinearity that are observed when optical field intensities become large. Resonant structures, such as high-Q ring microresonators and photonic band gap (PBG) device structures, and reduced dimension silicon photonic waveguides can amplify optical field intensities (e.g., 3106 V=m for 1 mW optical power propagating in a 100500 nm silicon photonic waveguide) to the point that these fields produce nonlinear charge displacements, resulting in SHG, optical rectification, and all-optical modulation phenomena. Practical applications of organic electro-optic materials depend on the fabrication of thin films (100 nm to several millimeter thicknesses depending on various applications with film thicknesses of 1–3 mm required for most applications). Electro-optic films must also be characterized by large N and hcos3 ui. Realization of films of high optical quality (optical propagation loss of 2 dB=cm or less) requires careful control of the material’s chemical and physical properties, such as the content of H-containing polar bonds (e.g., OH, NH) and solubility in spin-casting solvents. Frequently, modest glass transition temperatures will be desired for intermediate processing stages such as electric field poling or nanoimprint lithography but a very high glass transition temperature will be desired for the final material to maximize thermal and photochemical stability of devices. Reversible Diels–Alder chemistry [28] provides a convenient way of tuning glass transition temperature whereas additional lattice hardness can be achieved utilizing irreversible crosslinking reactions such as those based on the thermally initiated free radical reaction of the fluorovinyl ether moiety [29]. Telcordia standards require long-term stability at 858C. A material glass transition of 1508C normally will permit Telcordia requirements to be surpassed. The highest glass transition temperature that we have achieved for an organic electro-optic material is of the order of 2008C. This value, which was realized with some considerable difficulty, was achieved without compromise of electro-optic activity or optical loss [27,30,31]. Realization of a final glass transition temperature of 2008C requires processing the material above 2008C, which in turn requires the chromophores to exhibit very high thermal stability (250–3508C). Organic electro-optic materials, like organic light-emitting device (OLED) materials, must exhibit adequate thermal and photochemical stability. The demands are less severe for electro-optic material as current flow is not involved in electro-optic devices and optical operating wavelengths are typically telecommunication wavelengths (1300 or 1550 nm bands), which are far removed from the interband electronic transition. Singlet oxygen chemistry is the most problematic photodegradation mechanism for organic materials and chromophore interband (HOMO–LUMO) absorption plays a critical role in singlet oxygen activation. That is, photostability is observed to depend on the separation between the operational wavelength of the device and the charge transfer absorption maximum, lmax. Indeed, a photostability figure-of-merit, B=s, can be defined where B1 is the probability of photodecay from the LUMO charge transfer state and s is the interband (charge transfer) absorption coefficient [32–37]. Packaging of electro-optic devices to minimize oxygen exposure appears one of the most straightforward ways of achieving long-term operation at telecommunication power levels (e.g., 20 mW). Since organic electro-optic devices must generally be integrated with other disparate material technologies in operational systems, integration and device engineering of organic materials is a critical issue in their commercial application. Electro-optic materials must, for example, be surrounded by cladding materials to prevent propagating telecommunications light from seeing ‘‘lossy’’ metal drive electrodes. Organic electro-optic waveguides must also be interfaced with low-loss silica waveguides in a manner that minimizes coupling loss. Crystalline lithium niobate electro-optic and semiconductor (GaAs and InP) electro-absorptive materials face similar challenges of integration but the individual considerations for these disparate materials are quite different. It is interesting to note that the cost of packaging materials (rather than materials costs themselves) will almost certainly dominate the cost of commercial
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devices. Organic electro-optic materials afford some impressive advantages relative to crystalline and semiconductor materials. Flexible and conformal devices can be fabricated by lift-off techniques [38,39]. Organic electro-optic materials can be directly integrated with both silicon electronics and photonics, opening the way for high-density ‘‘opto-chip (integrated photonic=electronic)’’ devices that may be critical to next generation computer chip technology and to airborne information management platforms [7–9,25,40,41]. Organic electro-optic materials are amenable to soft and nanoimprint lithography, opening the door to the possibility of cost-effective mass production of complex circuitry [42– 44]. Indeed, organic electro-optic materials appear particularly well suited for the production of a wide array of devices based on novel stripline, ring microresonator, PBG, cascaded prism, Fabry–Perot etalon, and superprism architectures. The major role of organic electro-optic materials may be the facilitation of new applications rather than simply competing with current materials, such as lithium niobate, for sales of single Mach Zehnder modulators. Organic electro-optic materials are in a very exciting stage of development. In the past 12 months, the electro-optic activity of materials has been more than doubled, reaching values in excess of 300 pm=V. New design paradigms have been identified that suggest that values in excess of 1000 pm=V may be achievable. A wide variety of new device structures have been demonstrated, as have unique processing advantages of organic electro-optic materials. Soft and nanoimprint lithography have been used to fabricate sophisticated coupled ring microresonator structures that exhibit exceptional device properties with individual ring diameters smaller than the diameter of a human hair. Organic electro-optic materials have been successfully integrated with silicon photonic circuitry, leading to demonstration of record high bandwidth optical switching and optical rectification. Good thermal and photochemical stability has been demonstrated as has stability in the presence of space radiation (gamma rays, high-energy protons) [45]. In this chapter, we attempt to provide an appropriate perspective from which to view these transformative events and to discuss the factors that ultimately will limit the improvement of organic electro-optic materials. For some properties such as electro-optic activity, excellent guidance for the development of improved materials is provided by theoretical (quantum and statistical mechanical) methods. For other properties, such as optical loss (which are no less importance), the process of developing theoretical methods for guidance in the systematic improvement of materials is only a beginning. For still other properties, such as thermal, photochemical, and radiation stability, material improvement is largely Edisonian in nature depending solely on chemical knowledge and intuition. This chapter also emphasizes the importance of auxiliary materials, such as cladding materials, in defining device performance. The virtually infinite variability of organic electro-optic materials has been cited as an advantage but without rational guidance for the systematic improvement of materials the process becomes prohibitively time-consuming. This chapter attempts to provide a perspective not only on recent advances but also on remaining challenges.
6.2
Theoretically Inspired Optimization of Molecular First Hyperpolarizability
Quantum mechanical calculations have been used for several decades to guide the improvement of molecular hyperpolarizability [1–10 and references contained therein]. However, with the exception of predicting the increase of hyperpolarizability with the length of the conjugated p-electron system of a chromophore, the application of theory has been a mixed bag of success and failure. In some cases, comparison of theory and experiment has been a matter of comparing apples and oranges. Theoretical calculations are typically carried out on single isolated molecules compared to experimental measurements where chromophores are typically dissolved in some solvent (inert host matrix). Theoretical calculations are normally performed in the long wavelength (zero frequency) limit whereas experimental measurements are typically carried out at wavelengths between 1 and 2 mm. Moreover, multiphoton effects and chromophore intermolecular electrostatic interactions can complicate experimental measurements. Also, the experimental conformation of the molecule may be different from the conformation on which theoretical calculations were executed. Molecular first hyperpolarizability can vary
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dramatically with twist of the p-electron system or disruption of bond length alternation in an extended p-conjugated system. Also, in some cases, there are a number of conformations that are of approximately the same energy and no significant energy barrier to conformational interconversion exists. In such cases, the molecular hyperpolarizability may reflect the averaging of contributions from several conformations. Molecular aggregation can also occur in some cases with the consequence of dramatically influencing measurement of molecular first hyperpolarizability. An ongoing problem with theoretical calculations is that it is difficult to apply more rigorous computational methods to the largest chromophores and trends that apply for shorter model compounds may or may not hold for longer versions. Of course, in practice, these larger and more complicated molecules exhibit the largest electrooptic activities and are of most interest for device prototyping. Attempts at using theory to guide the improvement of molecular first hyperpolarizability can be divided into two classes: (1) A modular approach to improvement of the basic donor–bridge–acceptor structures that focuses on identifying the best donors, bridges, and acceptors [46] and (2) exploration of novel chromophore architectures such as ‘‘X’’-shaped chromophores [47–49] and twisted chromophores [50]. Note that it is difficult to define absolute values of b for a number of reasons and researchers attempting to identify and synthesize new improved molecules normally focus on relative b values. The utility of theory in predicting trends with simple variation bridge and acceptor moieties is illustrated in Table 6.1 [46]. In a number of cases, Hartree–Fock (HF), density functional theory (DFT), Green’s function theory, and semiempirical INDO calculations agree on trends [51] (see Figure 6.2). In other more complicated cases (e.g., twisted chromophores), they do not. In some cases, large predicted improvements in molecular hyperpolarizability are not realized in experimental practice. An example is the predicted improvement in b with the use of ‘‘gradient bridge’’ chromophores [52]. A gradient bridge chromophore has the general structure strong donor–weak TABLE 6.1 Comparison of Experimental and Calculated βHRS Values Relative to p-nitroaniline for Four Chromophores
F3C
O Fe
O CN
NC
CN
Fe
NC
NC
1
NC
2 NC
NC
CN
NC O
O
S
S
3 r33(pm/V) 1.3 µ@20%
F3C
Fe
Fe
Cmpd #
CN
NC
4 Experimental b relative to pNA
DFT Calculations
calc b rel,zzz
1
—
3.5
4.6
2.4
2
5
5.7
4.8
3.6
4/3
25/—
42.2/33.3
43.5/35.5
44/11.4
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b
103
102
PBE INDO HF B3LYP
101 101
102 b (PBE)
103
FIGURE 6.2 Benchmarking of Hartree–Fock (HF), intermediate neglect of differential overlap (INDO) semiempirical, and DFT PBE and B3LYP calculations have been carried out by Bruce Robinson, Jean-Luc Bredas, and Andy Chafin for 12 common chromophores differing in choice of donor, bridge, and acceptor moieties. Molecular first hyperpolarizability, b, values relative to p-nitroaniline have been calculated for each method and values obtained by HF, INDO, and DFT B3LYP are plotted against values obtained using DFT PBE (straight line in the figure). Proximity to the straight line illustrates agreement between the methods in predicting simple structure–function trends for b.
acceptor–strong acceptor. An example is a weak donor–weak acceptor bridge consisting of thiophene– thiazole units, which is theoretically predicted to exhibit a larger b than an analogous thiophene– thiophene bridge [52]. The experimental b values (determined by hyper-Rayleigh scattering [HRS]) suggest that only modest improvement in b is realized; these results appear to be more consistent with more recent results of DFT calculations [53]. An example of a trend that is predicted by various types of quantum mechanical calculations and that holds for both long and short bridge involves replacement of the furan ring of CLD-type acceptors [54] with pyrroline and pyrrolizine ring systems. Such replacement is predicted to lead to significant (several hundred percent) improvement in molecular hyperpolarizability and appears to be borne out by experimental measurements [44]. Sometimes a structure–function trend that is correctly predicted for shorter model compounds does not translate to larger (longer) chromophores where theoretical calculations are more difficult. An example is the dependence of b (and r33) on variation of donor structure in going from an amine donor to a guanidine donor. When complex architectures are examined, the situation becomes both more interesting and more complex than with simple donor, bridge, and acceptor variation. For ‘‘X’’-shaped chromophores, theory correctly predicts a ‘‘blue-shift’’ in the lmax of the charge transfer band and a modest increase in b. However, aggregation of these chromophores occurs in practice contributing exciton bands and complicating the utilization of such chromophores. Further structural modification (nanoscopic engineering) of these chromophores will be required before they can be unambiguously characterized and used for device applications. Twisted chromophores [50] are even more complex (see Figure 6.3 and Figure 6.4) and various theoretical methods do not agree. This divergence in predictions can be partially attributed to singularities and to parameterization inherent in semiempirical methods but clearly theory does not provide
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N
O
+ N
O
N
O
Semi-empirical hyperpolarizability vs twist angle 500
bm(10−30 esu)
0 −500
30
40
70
50
80
60
90
−1000 −1500 −2000 −2500 −3000
Twist angle (degrees)
−3500 MRD-CI INDO/S
CIS INDO/S
FIGURE 6.3 Molecular first hyperpolarizability, b, values calculated by semi-empirical INDO methods are shown as a function of twist angle for a ‘‘twisted’’ chromophore (see upper part of the figure). The maximum predicted value of b is large and negative, suggesting a zwitterionic ground state.
high confidence for pursuing twisted chromophores. However, these materials should be investigated if for no other reason than for the insights that such an investigation would provide into the utility of various theoretical methods for investigating this class of materials. In summary, quantum mechanical calculations can provide importance guidance in identifying chromophore structures, leading to improved molecular hyperpolarizability; however, such guidance is not foolproof and should be validated by preparation of more easily synthesized model compounds before launching into the synthesis of complicated chromophore molecules with all of the requirements necessary for utilization in device prototyping.
Ab initio hyperpolarizability for twisted mercyanine day 800
b (10−30 esu)
600 400 200 0
30
40
−200
50
60
70
80
90
Twist angle (degrees)
−400 B3LYP/6-31+G** HF/6-31+G**
(14,13) CASSCF/6-31+G* MP2/6-31+G**
BS-UBLYP/6-31+G**
FIGURE 6.4 Molecular first hyperpolarizability, b, values calculated by various ab initio methods are shown as a function of twist angle for the same chromophore as in Figure 6.3. In some cases, the maximum b value is positive, suggesting a singlet biradical ground state.
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l / 2 plate
Focusing lens Close-up of sample-holder Sample with flow cell
Polarizing cube Detection Computer
Integrating Focusing cavity lens
LED
Laser spectrum analyzer
Camera
Pump laser
TI:sapphire oscillator
Optical parametric oscillator
Oscilloscope
FIGURE 6.5 A schematic representation of a femtosecond, wavelength-agile (700–2000 nm) hyper-Rayleigh scattering (HRS) apparatus capable of absolute (integrating sphere) b measurements is shown.
6.3
Characterization of Molecular First Hyperpolarizability
Molecular first hyperpolarizability is typically estimated or measured using solvatochromism, electric field-induced second harmonic generation (EFISH) [55], or HRS [44,56] measurements. The relationship between solvatochromism (the shift in absorption maximum, lmax, with changing solvent dielectric constant) and molecular first hyperpolarizability can be understood as the effect of changing solvent electric fields, altering the contributions of the neutral and charge-separated forms to the ground and excited electronic states. The solvent fields mimic the externally applied field. However, solvatochromism measurements are not reliable indicators of true structure–function relationships relevant to optimizing molecular first hyperpolarizability and use of such measurements for estimating b has essentially stopped. Thus, HRS and EFISH measurements have become the defined standards for characterization of molecular first hyperpolarizability. Even these measurements are not without difficulties. Indeed, nearly an order of magnitude uncertainty exists in b values for common materials such as chloroform. This is quite worrisome particularly when it is realized that materials such as chloroform are commonly used as reference standards for determining b values for other materials. A number of problems contribute to uncertainty in the measurement of b, including two-photon absorption and fluorescence. Also, aggregation of chromophores can complicate measurements. Measurements carried out over a range of wavelengths are typically required to understand the contribution of resonance enhancement to molecular first hyperpolarizability. Such knowledge is, in turn, critical for a comparison of molecular first hyperpolarizability among different types of chromophores. Measurements as a function of chromophore concentration in one or more solvents are also typically required. A femtosecond-time-response, wavelength-agile HRS apparatus being used with some success [44] is shown in Figure 6.5. Also shown is a modification of the apparatus for the measurement of absolute b values. This instrument has been engineered using the latest technology for high signal-to-noise measurement capability as well as wavelength agility. Solution concentrations as low as nanomolar have yielded useful signal-to-noise.
6.4
Synthesis of Chromophores
The synthesis of a typical electro-optic chromophore is shown in Scheme 6.1. The synthesis of multichromophore dendrimers is shown in Scheme 6.2 and Scheme 6.3. Although the discussion of the synthesis of the multitude of chromophores that have been prepared to the present is beyond the scope of this chapter, it should be noted that optimization of reaction yields is an important requirement. In this
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Conjugated Polymers: Processing and Applications NaBH4 S
S
NaOH
O
MeOH Br
1)
OTBDMS O N
HCI(conc)
HO
quant 2)
14) KO(t-bu)
OTBDMS
7)
S 6)
TBDMSO
NaBH4 NaOH MeOH/THF 0°C-RT 79%
N
HO CN
OTBDMS N
Et2O −78°C 90%
S 7) O
OTBDMS N 9) HO NC
NC
OTBDMS
O
S
THF −78°C-RT 87%
8)
NC
Br
n-Buli DMF
S
NC
9)
n-Buli (1-equiv.) DMF
N
5)
P(OEt)3
O P O heat 3 days O Br quant 3)
Br
0°C-RT quant
THF
S
CI
CN
O Ph CF3
S
N
132 = ETOH RT 12 hr 62%
O CF3
D2.pas.27 HO
SCHEME 6.1 O OH O
O
10) O
HO O TBDMSO N
NC NC CN S O CF3
DCC/DPTS CH2CI2 50°C 24 hrs 40%
HO D2.pas.27 CN CF NC O 3 NC O O S
N
O
O
O O
CN O CN CF3 CN
O
O O F3C O S CN NC CN D2.pas.33
SCHEME 6.2
OTBDMS
S
N OTBDMS
O
N OTBDMS
OH
S
4)
Br
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NC
N
CN O CF
OTBDMS
NC S
O O
O
O
O O
S CN O CN CF1 CN
O
N
F
F
O
F
OTBDMS O O F1C O CN NC CN
S
N
HCI acetone quant
F
F
F
O F
F
F F F
O OH
F
F
DCC/DPTS CH2Cl2 62% CN CF1 NC
OTBDMS
D2.pas.41
N
F F O F
O O
O F
O
D2.pas.33
F
F F
NC S F F
O O
O
O
O O
S CN CN O CF1 CN
F
F
O O
O
F
N
O F
F
F F
F
O
O O F1C O CN NC CN
N
S
O
O
F O
F
O
F
F F
F
F
F F F
SCHEME 6.3
regard, microwave-assisted synthesis [57] has become an important tool in the arsenal of the synthetic organic chemist.
6.5
Theoretically Inspired Optimization of Macroscopic Electro-Optic Activity
6.5.1 Chromophore–Polymer Composite Materials By chromophore–polymer composite materials, we refer to chromophores physically incorporated (dissolved) into commercially available polymer materials such as amorphous polycarbonate (APC) [58] from Aldrich Chemicals. Chromophore and polymer are dissolved in a suitable spin-casting solvent, such as cyclopentanone. Spin-cast thin films are heated to near the glass transition temperature of the composite material (which will vary with chromophore concentration due to the plasticizing effect of the chromophore). Acentric chromophore order is induced by electric field poling. If one assumes that the presence of the polymer host does not sterically hinder the reorientation of the chromophores under the influence of the poling field, the order parameter can be readily calculated. We have already noted that if chromophore–chromophore intermolecular electrostatic interactions are neglected then hcos3 ui ¼ mF=5kT and the order parameter will be independent of chromophore concentration (or number density, N). Intermolecular electrostatic interactions can be treated at several levels of sophistication. If chromophores are treated as point dipoles (this amounts to neglecting the effect of nuclear repulsive interactions, i.e., size and shape of the chromophore), an analytical expression can be readily derived [59– 67] that illustrates that electronic intermolecular electrostatic interactions contribute an ‘‘attenuation factor’’ to the order parameter expression; namely hcos3 ui ¼ mF=5kT[1 L2 (W=kT )] where L is the Langevin function and W is the intermolecular electronic electrostatic interaction energy. Key to the derivation of this analytical expression is the computation of a potential energy function that describes the electric field at a selected reference chromophore contributed by the ensemble of surrounding chromophores. This potential function combines in a vectorial manner with the chromophore dipole moment–electric poling field interaction (mF ) potential to define the total electric field (potential energy) experienced by the reference chromophore. The integrals over orientational variables required to compute hcos3 ui are quite complex but the problem can be simplified by noting the trigonometric relationship that
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exists between the axes systems of the reference chromophore, the applied poling field, and the potential contributed by surrounding chromophores [64–67]. Nuclear repulsive interactions can be treated as r12 interaction potentials or as hard object potentials (the repulsive potential goes to infinity as the chromophores start to overlap). In the latter approximation, the effect of nuclear repulsive (or steric) interactions among the chromophores can be treated by adjusting the limits of integration of the integrals over angular (orientational) variables [64–67]. Computations based on the aforementioned simple treatment of electronic and nuclear intermolecular interactions among chromophores lead to a surprising good reproduction of experimental data without the introduction of adjustable parameters [59–67]. The relative magnitudes of the thermal energy (kT ), electric field poling energy (mF), and the intermolecular electrostatic energy (W ) will define the ratio r33=r13. If thermal energy dominates the other two, the ratio will be approximately 3. In the limit of our analytical derivation of order parameters treating electronic-only intermolecular electrostatic interaction in the point dipole approximation, increasing W acts to drive the ratio r33=r13 to 3 even when mF > kT. At low W and modest kT, the ratio will be greater than 3 for large mF with the exact ratio dependent on the ratio of mF to kT. If nuclear repulsive interactions are taken into account, the situation is more complex and rigorous treatment of all intermolecular electrostatic interactions suggests that the ratio can be either greater or smaller than 3 depending on chromophore shape. Oblate ellipsoidal chromophore shapes favor ratios greater than 3. The Ising lattice limit also yields ratios greater than 3 with the precise ratio depending on the anisotropy of the molecular first hyperpolarizability tensor. The dependence of electro-optic activity upon poling field strength becomes nonlinear at large F in the Ising lattice limit. Of course, one can also use Monte Carlo methods [62,68] to consider the effect of intermolecular electrostatic interactions among high dipole moment chromophores. The potential function (describing the field produced by surrounding chromophores on a reference chromophore) from Monte Carlo methods is very similar to the potential function derived analytically [62,65,67]. Monte Carlo calculations have the advantage of permitting a rigorous treatment of electronic and nuclear electrostatic interactions. The accuracy of such calculations typically relate to the size of the system treated (e.g., an ensemble consisting of 1000 chromophores). Results of Monte Carlo calculations of the loading parameter, N hcos3 ui ¼ (r33=b) (constant), as a function of chromophore number density N are shown in Figure 6.6. Curves for various chromophore shapes are given. Also shown in the figure are curves for the independent particle case and for chromophores experiencing an Ising potential. Note that the variation of nuclear repulsive interactions with changing chromophore shape acts to change the relative importance of components of the intermolecular electronic electrostatic potential function that respectively favor acentric versus centric chromophore order. Also, note that the highly confining Ising potential leads to very large electro-optic activity (that is, a very large r33; r13 is small in the Ising model). An extremely important observation derivative from this computational data set is that nanoscopic engineering of chromophore environment can lead to intermolecular electrostatic interactions acting to enhance poling-induced acentric order. Thus, order above the independent particle limit (mF=5kT ) can be achieved. Indeed, order parameters, hcos3 ui, in the range 0.5–1.0 are possible for finite poling fields. Self-consistent field theoretical methods [69–73] have also been applied to this problem with similar observations. Fundamentally, theory has done a respectable job of explaining variations of electro-optic activity with simple changes in chromophore shape. There are numerous examples in the literature where electro-optic activity has been increased by making originally prolate ellipsoidal chromophores more spherical or more discotic in shape. We show one example in the accompanying table (Table 6.2).
6.5.2
Dendrimers and Super=Supramolecular Materials=Organic Glasses
When chromophores are covalently coupled in super=supramolecular objects (such as multichromophore-containing dendrimers), Monte Carlo-molecular dynamic calculations must be modified to take into account the restrictions on motion associated with covalent bond potentials. To accurately account for covalent bond potentials, ‘‘atomistic’’ Monte Carlo methods are required [68]. However because of the large number of atoms involved, fully atomistic calculations would be prohibitively time-consuming
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x 1019 Region of enhanced order
18 Ising lattice
2:1 Oblate
Loading parameter
16 14 Independent particle lattice
12 10
Spherical
8 6 4 1:2 Prolate 2 0 0
0.5
1
1.5
2 2.5 3 3.5 Number density (molecules/cc)
4
4.5
5 x 1020
FIGURE 6.6 The variation of N hcos3 ui (loading parameter) with number density N is shown as a function of chromophore shape. All data were calculated by Monte Carlo methods [61]. Also shown are two limiting forms: The independent particle and Ising lattice limits. The shaded region between the straight lines corresponding to these limiting forms represent a region where intermolecular electrostatic interactions enhance poling-induced acentric order. Below the independent particle limiting line, intermolecular electrostatic interactions act to attenuate polinginduced acentric order. This behavior reflects the fact that there are two components to the intermolecular electronic interaction potential, one favoring centrosymmetric order and one favoring noncentrosymmetric order.
and expensive for the systems of the size that we are considering [62]. One answer to this dilemma is a ‘‘pseudoatomistic’’ approach. In this pseudoatomistic approach, flexible segments of the chromophorecontaining macromolecules are treated using fully atomistic methods with potential functions defined by quantum mechanics. Segments involving p-electron conjugation are treated in the united atom approximation (justified by the fact that p-conjugation restricts rotations about bonds). In all cases, electron distributions are derived from quantum mechanical calculations. An example of a ‘‘pseudoatomistic’’ computational structure is shown in the accompanying Figure 6.7 for a three-arm (with a chromophore in each arm attached in a ‘‘side-on’’ manner) dendrimer. For such multichromophorecontaining dendrimers, pseudoatomistic Monte Carlo calculations show that acentric order parameter, hcos3 ui, increases with chromophore number density (e.g., concentration of electro-optic dendrimer in a host polymer such as APC). Indeed, as is shown, in Figure 6.8, a factor of 3 improvement in electrooptic activity is observed for a pure dendrimer film relative to the best electro-optic activity achievable for the same chromophore in an APC composite. Monte Carlo calculations also predict that the ratio of r33=r13 will increase and indeed this ratio is another effective measure of order and the relative magnitudes of kT, mF, and inter (and intra)molecular electrostatic energies. Even more dramatic results are observed for binary organic electro-optic glasses as in shown in Figure 6.9. By ‘‘binary’’ we refer to the fact that the glass consists of two different types of chromophores. Moreover, no inert host polymer is involved. Obviously, chromophore loading (number density) is quite high for such glasses. The experimentally observed, large electro-optic coefficients indicate that the acentric order parameters are quite large. How can this surprising and important result be understood? First of all, we note that the situation can be even more complex than the cases discussed to this point. Van der Waals interactions are very
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TABLE 6.2 Relative r 33 Values Are Shown for Four Chromophoses Dissolved in APC OMe
N
BU
BU
O
O
Me O N
S
BU
NC
BU
O
S CN
O
O
O
S
O
S
CN
O
O
NC
NC
Me
NC
OMe
OLD-1
OLD-2
OMe
Me N
S
OLD-3
O
N
S
O
S
O
S CN
NC NC
O
F3C
O CN
O NC
Me
NC
OMe
OLD-4
Wt%
20%
30%
OLD-1
3.0
—
OLD-2
2.0
3.6
OLD-3
1.0
—
OLD-4
4.0
—
important in these binary glasses and this includes, in some cases, interactions such as H–F hydrogen bonding. It is very easy, and potentially problematic, to oversimplify the situation but we note that a somewhat intriguing (and overly simplistic) interpretation is to consider such glasses as approximating an Ising lattice where ordered chromophores impose a restrictive (steric) potential on other chromophores.
FIGURE 6.7 The pseudo-atomistic Monte Carlo representation of the D2.PAS.41 three-arm dendrimer of Scheme 6.3 is shown.
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r33 at 1310 nm wavelength (pm/V)
200 CF3FTC D2PAS
150
100
50
0 0
0.4
0.2
1 1.2 0.6 0.8 Poling field (MV/cm) NC
1.6
NC
TBDMSO
CN S
N TBDMSO
1.4
O CF3
CF3-FTC
FIGURE 6.8 Electro-optic activity (measured at 1.3 mm) as a function of poling voltage is compared for a chromophore (CF3-FTC)=APC composite and for the D2.PAS.41 dendrimer of Scheme 6.3. The composition of the (CF3-FTC)=APC composite was adjusted to yield maximum electro-optic activity. The data shown are, in each case, a composite of two preparations. The chromophore is the same for both preparations so the differences can be attributed to differences in intermolecular electrostatic interactions.
400
O O
O
350
N
O
O NC
CF
F
F F
2
F
F
O
NC
F
F
250 200
F
O
O
S
r33 (pm /V)
F
O
300
F CN
150 Si
O
N
O
Si
100 50 0 0
50
100
Poling voltage (V/µm)
150
CF2
NC O
NC
S CN
FIGURE 6.9 Electro-optic activity (measured at 1.3 mm) for chromophore-doped, chromophore-containing dendrimer glasses is compared to that of the undoped D2.PAS.41 dendrimer (denoted by crosses) of Figure 6.8. Data are plotted as a function of poling field. The circles and best-fit solid line denote data for the D2.PAS.41 dendrimer doped with a CLD-like (see chromophore in lower right for general structure) chromophore. Also shown are data for various concentrations of CLD-like chromophores doped into the single-chromophore-dendrimer shown in the upper right of the figure.
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Binary organic glasses yield electro-optic coefficients in the range 300–400 pm=V (more than an order of magnitude greater than lithium niobate) even when chromophores with modest molecular first hyperpolarizability values are employed. Chromophores with pyrroline- or pyrrolizine-ring (rather than furan-ring) acceptors should lead to electro-optic activities greater than 600 pm=V. It is interesting to speculate (from the perspective of current theoretical understanding discussed here) what ultimate electro-optic activity can be achieved. The first observation that can be made is that the loading parameter, N hcos3 ui, cannot be significantly improved from current values. A factor of less than 3 is likely the upper limit for future improvement. Thus, improvement in electro-optic activity must, in the future, largely come from improvements in b. Clearly, molecules with larger b values are likely to be conceived and synthesized. However, for such molecules to lead to practical electro-optical materials, optical absorption loss must be taken into account. For ‘‘neutral’’ ground state (or positive b) chromophores, the interband absorption maximum typically shifts to longer wavelengths with increasing b. The relationship is not as simple as has been suggested (a linear relationship with a universal slope) but increasing b significantly beyond the values for molecules considered here is likely to lead to detectable optical loss (e.g., greater than 1–2 dB=cm) at 1.3 mm telecommunication band wavelengths. Although larger electro-optic activity permits shorter device lengths to be employed and thus material loss is not such a critical issue (see subsequent discussion in this chapter), a higher b value does not provide a significant advantage if the ratio of b to optical loss does not improve. Indeed, the ratio of r33 to optical loss is an important component of a material figure of merit for many device structures. When optical loss considerations are taken into account, it is unlikely that useful electro-optic coefficients will extend much beyond 1200 pm=V and even realization of 1200 pm=V will be a major challenge. Of course, the relative importance of material loss varies from one class of devices to the next. For example, material optical loss is not such a critical issue for ring microresonator devices (including silicon photonic hybrid devices) because of the high bending losses associated with such devices. It should also be kept in mind that while highly ordered materials lead to large r33, such order does not enhance r13. Thus, device structures (such as certain etalon structures) that depend on utilization of r13 may not benefit from improvement in acentric order hcos3 ui.
6.6
Characterization of Electro-Optic Activity
A variety of techniques including simple refection or Teng–Man measurements [74], attenuated total reflection or ATR measurements [75], ellipsometric measurements [76], two-slit interference measurements [77], and Mach Zehnder interferometric measurements [78] have been routinely used to characterized electro-optic activity. An important component of characterization of electro-optic activity has been in situ monitoring of the introduction of electro-optic activity by electric field poling and the subsequent relaxation of electro-optic activity after the poling field has been turned off. In such measurements, relative electro-optic activity (or actually acentric order hcos3 u i) is the quantity of interest and SHG [26] has frequently been employed. However, absorption of the second harmonic light (as it excites the chromophore interband charge transfer absorption) can be a problem and more recently in situ monitoring has relied on modification of direct electro-optic measurement techniques such as the Teng–Man method (shown in Figure 6.10). In the modified Teng–Man method, a heating and temperature measurement stage has been added as to have capabilities for applying a DC bias (i.e., poling) field. Examples of monitoring the introduction and relaxation of electro-optic activity as a function of thermal ramping are shown in Figure 6.10 (insets A and B). Such methods provide a practical means of defining material glass transition temperature and yield numbers in good agreement with traditional differential scanning calorimetry (DSC) measurements. Instruments are also frequently modified to permit measurement of conductivity as a function of temperature. Such measurements provide insight into the actual magnitude of the electric poling field felt by the electro-optic material and provide yet another method of defining glass transition temperature. The modified Teng–Man apparatus is also useful in defining the activation barrier to relaxation of poling-induced electro-optic activity. By measuring the temporal variation of electro-optic activity as a function of temperature for a range of
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1000⫻(Im/Ic)
3
A 2
1
V (volt)
0 55
Laser λ=1310nm
V (t)=Vd.c + Vo sinωt
60
65
70 75 80 85 90 Temperature(°C) 3
95
Detector 1000⫻(Im/Ic)
Time Lens Polarizer
Polarizer
Lens Lens
2 B 1
SB Compensator 0 30 40 50 60 70 80 90 100 Temperature(°C)
Heating block Thermocouple Heater controller
Heater
I-V converter
ITO Oscilloscope Gold
Polymer film
Nitrogen atmosphere Thermometer keithleyn DMM2000
Lock-in amplifier
Ampere meter keithley2400
DC power supply
Function generator
FIGURE 6.10 Modification of a TengMan apparatus [73] to permit in situ monitoring of induction of electro-optic activity by the poling field and subsequent relaxation of that activity after the poling field is turned off is shown. This apparatus permits simultaneous application of a large DC bias (poling) field and a high-frequency field for assessing electro-optic activity. Inset A shows data recorded for poling a multichromophore-containing dendrimer whereas inset B shows relaxation of the electro-optic activity after the poling field is turned off. To obtain the data traces shown in A and B, the temperature was increased at a rate of 108C=min while monitoring the Teng–Man signal.
elevated temperatures and by employing stretched exponential fitting of the kinetic data, activation barriers and material thermal stability can be defined as shown in Table 6.3. In this table, we compare the thermal stability of two multichromophore-containing dendrimers differing only in mode of attachment of the chromophores to the dendrimer cores. Side-on attachment leads both to higher electro-optic activity and to greater thermal stability of that activity. A similar type of modification [78] can be applied to the commercially relevant Mach Zehnder interferometer device. This is a more complicated undertaking as appropriate cladding materials must be identified. However, this test bed configuration has the advantage of not only permitting material properties and processing conditions to be studied in situ but also permitting the systematic exploration of various device designs through the relationship between r33 and drive voltage (Vp, the voltage required to produce a phase shift of p or equivalently to turn the signal from the interferometer ‘‘on’’ and ‘‘off ’’), e.g., Vp ¼ lh=n3 r33LG where l is the operating wavelength, h is the electrode spacing, L is the electrode length, and G is the modal overlap parameter. A factor of 2 must be added to the denominator if push–pull operation is used.
6.7
Auxiliary Properties and Their Characterization
Optical loss can be divided into three categories: (1) material absorption loss, (2) material propagation loss (the sum of absorption and scattering loss), and (3) total electro-optic device insertion loss (the sum of absorption, scattering, and coupling loss). Typically, coupling loss will dominate the total insertion
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TABLE 6.3 Comparison of Thermal Stability NC NC
N
CN
O
S O
S
CN
O
O
O
CN CN
NC
O
O
O
O
NC
O
O
N
O
O
N
O
O NC
N
CN O
O O
O
NC
O
NC
S
N
O S
O
S
S
O
O NC
NC
CN CN
CN CN
N
End-on AALD-1104 V
Side-on D2.PAS.31 IV
HE-End-on V
HE-Side-on IV
t(s)
(s)
b
t (s)
(s)
b
118°C
47.39
817.54
0.337
133.37
666.57
0.3506
114°C
67.66
0.300
454.55
14354.25
0.227
119°C
409.33
0.256
2327.21
6037700.72
0.150
Ea (kcal/mole)
32669.02 124.57
302.84
loss followed by propagation loss and finally by material absorption loss. However, it should be noted at the outset of this discussion that it is no trivial matter to separate absorption and scattering loss in practice. Moreover, it is important to note that any of the above contributions to loss can be unacceptably large. Indeed, a great deal of effort must be expended to keep various loss contributions to acceptable levels. We now discuss each these contributions to optical loss in turn. The dominant contributions to optical loss associated with absorption of light at telecommunication wavelengths are (1) hydrogen vibrational overtone absorptions, (2) interband (charge transfer) electronic absorption, and (3) weak charge transfer (intragap) states associated with aggregation, hydrogen bonding, and intermolecular donor–acceptor interactions. However, if contributions from this last category are present, the optical loss is frequently unacceptably high and the materials are not useful for device prototyping. Ideally, materials are engineered so that the first contribution dominates. This is normally accomplished by selecting chromophores where the lmax for the interband transition is far removed from telecommunication wavelengths, where intermolecular association effects are absent, and where inhomogeneous broadening of the interband transition is not high. For ‘‘neutral’’ ground state chromophores, increasing b often results in some red [bathochromic] shifting of the interband absorption. While a simple linear relationship exists between b and lmax for some systematic chromophore structural modifications (e.g., increasing the length of the p-electron bridge for polyene bridges), this is not always the case and the magnitude (and even sign) of the slope of the variation can change with different types of structural modifications. An example where b increases slightly and the interband absorption is blue-shifted (hypsochromic shift) is that of ‘‘X’’-shaped chromophores [48,49]. Clearly, significant interband absorption at telecommunication wavelengths or absorption within the gap are unacceptable and ultimately place design restrictions on the modification of chromophores for improved molecular first hyperpolarizability. Hydrogen vibration overtone absorption represents a fundamental problem and typically limits absorption loss in organic electro-optic materials to values
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Organic Electro-Optic Materials
of 1 dB=cm or somewhat greater. This contribution to absorption loss can only be reduced by decreasing hydrogen (proton) density. Fluorination is a route to reduction of proton density and has been used to achieve optical loss of less than 1 dB=cm for organic electro-optic materials. Partially fluorinated electro-optic dendrimers have exhibited optical loss as low as 0.2 dB=cm at C-band (e.g., centered around 1.55 mm) wavelengths. It should be remembered that telecommunication wavelengths were originally chosen because these are the wavelengths associated with lowest hydrogen vibrational overtone absorption and engineering of ‘‘long-haul’’ silica fibers required minimization of hydrogen content. Photothermal deflection spectroscopy (PDS), widely used by researchers at Lockheed Martin [79,80], is an attractive way of investigating various contributions to ‘‘absorption’’ loss, although it is not clear if the PDS measurements are unaffected by scattering and are most useful in measuring the ‘‘loss spectrum’’ rather than absolute values of absorption loss. Lockheed Martin researchers have focused in the past several years upon attempting to bring a theoretical understanding of how various solid state interactions contribute to absorption loss measured by PDS methods. In developing new materials, it should be kept in mind that if absorption loss increases at the same rate as the increase in b, no gain in material performance has been achieved if the device interaction length is constant. This will not be the case if hydrogen vibrational overtone absorption dominates absorption loss; in this case, increase in b will be independent of absorption loss as long as interband electronic absorption does not contribute in a noticeable way. In this latter case, the issue of absorption loss can be addressed by simply using a shorter interaction length in the device structure if exceptional electrooptic activity exists. An example is shown in Figure 6.11 where a total insertion loss of less than 3 dB together with a drive (Vp) voltage of less than 1 V is predicted for a material with a material propagation loss of 2 dB=cm (which is currently viewed at the highest acceptable value). Short device lengths are also useful for increasing operational bandwidth due to reduction of resistive electrical signal losses in metal electrodes. In overview, many other sources of optical loss (scattering loss, 4
300
3.5
Material and taper parmeters:
3 200
2.5 2
150
1.5
100
Vπ (Volts)
Frequency (GHz)
250
1 50
d (thickness) = 8 mm m (wave loss) = 0.75 dB(GHz)1/2/cm Fiber coupling = 0.8 dB/coupling
0.5
0
Material loss = 2 dB/cm
0 0.1
0
0.2
0.3 0.4 0.5 0.6 0.7 Interaction length, L (cm)
0.8
0.9
1
4
4
3.5 3 2.5
3.5
Device parameters
3
L = 5 mm
2.5
2
2
1.5
1.5
1
1
Vπ (Volts)
Fiber to insertion loss (dB)
r 33 = 300 pm/V
Vp = 0.75V BW(3 dBe) = 90 GHz Insertion loss = 2.6 dB
0.5
0.5
0
0 0
0.1 0.2 0.3 0.4 0.5 0.6 0.7 0.8 0.9 Interaction length, L (cm)
1
FIGURE 6.11 The relationship of material and device parameters is shown for recently prepared organic electrooptic materials.
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coupling loss, bending loss in the case of ring resonator devices, etc.) frequently dominate the total device insertion loss and absorption loss is not an dominant issue that practically impacts performance, provided lmax of the interband transition is not permitted to closely approach operational wavelengths. Of course, what constitutes an acceptable lmax must be defined for each new electro-optic material and will depend on many features of the material system. The reader’s attention is drawn to a recent analysis of the relationship of absorption loss to the structural features of chromophore– polymer composite materials carried out by researchers at Lockheed Martin [79,80]. This correlated theoretical–experimental investigation permits rational design of materials to stay within acceptable loss limits. This is very important in that many aspects of the interaction of chromophores with each other in the solid state and with the surrounding host matrix can lead to unacceptably high absorption loss values and a great deal of time can be wasted by simply relying on trial and error experimentation to identify materials with acceptable absorption loss. Material propagation loss is comprised of two components: (1) Absorption loss discussed above and (2) scattering loss associated with material inhomogeneity and surface roughness. Scattering loss contributions can range from insignificant to unacceptably high and their magnitudes are normally associated with processing conditions. The processing steps that can lead to the introduction of significant scattering loss include: (1) spin-casting or other film deposition processes, (2) electric field poling and lattice hardening, (3) waveguide fabrication, (4) deposition of cladding layers, and (5) metal electrode deposition. The two most common methods of characterizing propagation optical loss are the out-coupling method of Teng [81] and the cut-off method. The former is nondestructive whereas the latter is destructive. As the name implies, the cut-off method simply involves measuring propagation through the material as a function of simply cutting off various lengths of materials. The method of Teng [81] has the advantage of permitting a series of measurements on the same sample with minimum perturbation of the sample and the measurement set up. Spin-casting is the most common method of preparation of thin films of desired 1–3 mm thickness. Material solubility in traditional spin-casting solvents, solution viscosity, and spinning speed are important variables that define material homogeneity and thus scattering loss. For multicomponent composite materials, care must be exercised to avoid phase separation. Typical scattering loss values should be a few tenths of a dB=cm but losses of many dB=cm can be observed if the aforementioned conditions are not controlled. Electric field poling and lattice hardening can introduce loss in several different ways for different types of material preparation. Phase-separation during electric field poling and lattice hardening can lead to very high scattering losses [7–9,82]. Electrophoretic migration of chromophores can occur unless inhibited by covalent bonds attaching the chromophore to the higher molecular weight host lattice. Imbalance in the stoichiometry of crosslinking reactions, out-gassing of reaction products such as water in lattice hardening chemistries (e.g., urethane crosslinking), and lattice perturbation by crosslinking coupling can lead to high scattering loss [7–9,82]. In general, the conditions of electric field poling and lattice hardening require careful definition and control to simultaneously optimize electro-optic activity and to control scattering losses. Another way optical loss can be introduced during electric field poling is by material damage. Such damage will be problematic near the material surface. Obviously, dielectric breakdown of the electro-optic material must be avoided. The very rough surface of indium tin oxide (ITO) electrodes can be problematic as electric fields increase at the tips of ITO spires. It should also be noted that poling under conditions of in situ optical monitoring of poling efficiency can represent very material-unfriendly conditions (high electric fields, potential for some current flow, high light intensity, high oxygen content, and a lattice that permits facile oxygen diffusion, etc.). Any of the abovementioned effects can lead to optical loss contributions of many dB=cm. In overview, when optimized conditions are carefully defined, materials can be repeatedly prepared where processing-defined scattering losses are a few tenths of a dB=cm. Surface roughness of electro-optic waveguides is a problematic source scattering loss in a variety of device structures ranging from stripline devices to ring microresonators. Surface roughness is most frequently introduced during reactive ion etching (RIE) of electro-optic waveguides or during
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deposition of cladding layers. In RIE fabrication of waveguides, three factors typically contribute to wall roughness and associated scattering loss: (1) The quality of the etch mask, (2) the RIE conditions (whether conditions of a chemical or physical etch apply), and (3) the hardness of the material. When everything is optimized, excess loss associated with RIE processing can be insignificantly small (e.g., 0.01 dB=cm [7–9,83,84]). The crucial factor appears to be to avoid ‘‘pitting’’ due to reactive ions having too much kinetic energy (a physical rather than a chemical etch) and due to soft materials rearranging after etching. Typically, oxygen contributes the reactive ions in etching of organic waveguides although CF4 with spin-on-glass techniques has been used to make deep vias to interconnect metal electrodes and drive electronics through planarizing layers of polymers in the vertical integration of electro-optic photonics with semiconductor VLSI electronics [7–9]. Lattice hardness is also the crucial factor in preventing pitting by solvents used to deposit upper cladding layers by spin-casting. To achieve acceptable loss values in depositing cladding layers, the electro-optic material layers must be essentially insoluble in the spin-casting solvent. The resolution of etch masks can be an issue particularly in the fabrication of small-ring microresonator structures. For ring microresonator fabrication, bending loss (defined by the index contrast between core and cladding materials and by the size of the ring) and surface roughness loss associated with mask resolution typically dominate propagation loss with intrinsic material propagation loss being less important than it is for stripline device structures. Cladding layers are used to confine optical fields within the active core of electro-optic waveguides and most particularly to prevent propagating evanescent optical waves from seeing lossy metal electrodes. Commercially available cladding materials typically exhibit optical loss values of several dB=cm and unless care is taken such materials can contribute to apparent waveguide propagation loss. Moreover, the thickness of cladding materials required to prevent optical fields from sensing metal electrodes will depend on the index of refraction contrast between the core and cladding materials. Unfortunately, thick cladding layers will attenuate the electric fields felt by the core electro-optic material and will translate into a requirement of a larger Vp for device operation. Thus, a happy compromise must be sought in cladding thickness. A potential route to the use of thinner cladding layers would be to employ ‘‘transparent’’ metal oxide electrodes. However, the poorer conductivity of current metal oxide materials can result in poorer bandwidth performance of devices, so high-purity gold electrodes continue to be the standard for high-bandwidth devices. It is also desirable for cladding materials to exhibit higher conductivity than core electro-optic materials so that the poling and radiofrequency (rf) [85–90] fields are dropped across the cladding materials, leading to better poling efficiencies and lower Vp operation. To this point in time, cladding materials with improved conductivities also exhibit increased optical loss negating their use. Indeed, one of the last great frontiers in organic electro-optic device technology may be the development of new cladding and electrode materials. Two factors define coupling loss between passive long haul silica fiber waveguides and active organic electro-optic waveguides: (1) index of refraction contrast between active and passive waveguides and (2) mode size mismatch. With lithium niobate electro-optic waveguides, the problem is the large index contrast whereas for organic electro-optic materials (with index of refraction values much closer to silica) the problem is mode size mismatch. An optical beam propagating in a 1.3 mm telecommunications silica fiber has a spherical mode size of approximately 10 mm. The optical mode in the organic electro-optic waveguide is an elliptical mode of typical dimensions 62 mm. If such waveguides are simply butt-coupled, a coupling (scattering) loss of 8 dB or greater will typically be observed. The answer to the problem of reducing coupling loss for organic electro-optic waveguides is to employ some type of mode transformer. A variety of structures have been proposed and demonstrated in the literature [7–9,91–96] and an example is shown in Figure 6.12. Spherical lens and grating coupling schemes have also been suggested but the former appears to afford poor mechanical stability and the latter has primarily found application in hybrid organic electro-optic=silicon photonic device structures. Key requirements for a successful mode transformer appear to be ease of fabrication and robustness. An issue with mode transformers is that the efficiency of coupling depends upon the length of the transformer structure. A reasonable target for such devices is a coupling loss of 0.7 dB. As suggested in Figure 6.11, a total insertion loss of 3 dB appears possible with new organic electro-optic materials while
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EO polymer waveguide Adiabatic tapers
Low loss large core waveguide
Output fiber core
Input fiber core
Measured mode patterns
Upper cladding not shown for clarity
Top electrodes
EO polymer waveguide
Chip loss-5.5 dB Fiber to waveguide Coupling loss ~0.45 dB
Passive polymer waveguide
ZPU-1215 ZPU-1279
FIGURE 6.12
Bottom electrode ZPU-1215 ZPU-1279
APC/CLD-1
An efficient mode transformer structure is shown together with performance parameters.
also exhibiting desired bandwidth and drive voltage performance. However, to achieve such performance, care must be given to minimizing loss at each processing and integration step and to minimizing instrinsic absorption loss. Indeed, controlling total insertion loss may be the single greatest challenge in developing functional electro-optic devices and systems based on organic electro-optic materials. There are two aspects to thermal stability: (1) Thermal decomposition temperature and (2) the temperature at which poling-induced noncentrosymmetric chromophore order is rapidly lost. Thermal gravimetric analysis (TGA) is usually employed to define decomposition temperature and typically only chromophores exhibiting decomposition temperatures above 2508C are carried forward in the development of practical electro-optic materials. The decomposition temperatures of chromophores (and of other components of electro-optic materials) define the ultimate temperature stability of electro-optic materials by defining the upper limit of poling and processing temperatures. It should be kept in mind that decomposition temperatures will vary with the state-of-matter (solid or fluid) and with the presence of oxygen. Because electric field poling is normally done in the presence of oxygen and involves conditions that maximize chemical decomposition, the thermo-chemical stability of chromophores typically restricts poling temperatures to 2008C or less. For electrically poled organic electro-optic materials, device lifetime is normally defined by the temperature at which poling-induced order is lost by rotational relaxation. The stability is related to difference between the glass transition temperature of the material and the device operating temperature. Normally, an operating temperature 508 below the material glass transition temperature will result in stability for many thousands of hours without significant (less than 5%) change in electro-optic activity. Thus, definition of thermal stability is, in the simplest sense, a matter of defining Tg. This can be done by DSC measurements or equivalently by the in situ Teng–Man and Mach Zehnder interferometric techniques discussed earlier in this chapter. To demonstrate that Telcordia Standards are satisfied, the device is heat-sinked at 858C and its operation is monitored as a function of time. For practical devices, electro-optic materials should exhibit glass transitions between 1358C and 2008C with the highest temperature materials required only for some
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specialized (largely military) applications. Engineering of high-Tg materials can be an expensive and time-consuming undertaking and should be pursued only when required by a particular device application or processing scheme (such as nanoimprint lithography). From a material structural perspective, thermal stability is defined by the number and placement of covalent bonds connecting components, by intermolecular electrostatic interactions, and by the segmental flexibility of components. In general, low glass transition materials are desirable for spin-casting and poling whereas high glass transition materials are desirable for latter stage processing and for stable device operation. Thus, it has become common practice to use crosslinking chemistries (see next section) to enhance thermal stability after the poling stage. Photochemical stability has been viewed by many as the Achilles’ heel of organic electro-optic materials [33–37]. In practice, photochemical stability of organic electro-optic materials can vary by many orders of magnitude [36,37] from unacceptable values to values that suggest that devices can be operated for decades without degradation in performance. Photo-decomposition of organic electrooptic materials at high optical powers appears to be most problematic in the presence of oxygen. Moreover, photostability appears to decrease with decreasing separation of the operating wavelength from the lmax of the interband (charge transfer) transition [33–37]. Photochemical stability measurements (see Figure 6.13) are most commonly effected by pump–probe measurements where the intensity of the interband transition is monitored as a function of pumping at the operational wavelengths. A problem with this approach is that intensity of the probe optical beam may cause photochemical decay rather than simply reporting such decay induced by the pump beam. This artifact will make the photoinstability of materials appear artificially large. This complication is made particularly problematic due to the fact that kinetic rates are slow even with the highest available (e.g., 1 W at telecommunication wavelengths) pump powers. A means of attenuating the perturbation from the probe beam is to sample the decay periodically (with the probe on only for short periods of time) while keeping the pump on constantly. To completely avoid the influence of an optical probe beam, it may be necessary to use an alternate detection scheme such as electro-optic monitoring. Until recently, few measurements of photochemical stability were conducted at telecommunication wavelengths. More recently, researchers at Corning [36,37] and in the Dalton research laboratory have carried out measurements at such wavelengths including over the C band (around 1.55 mm). With all measurements to the present, it has been common to report data in terms of a photochemical figure of merit, B=s, where B1 is the
FIGURE 6.13
A pump–probe apparatus for the measurement of photochemical stability is shown.
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Conjugated Polymers: Processing and Applications
probability that photochemical decay occurs from the LUMO state (excited charge transfer state) and s is the absorption coefficient of the interband electronic transition at the operational (pump) wavelength. For a typical chromophore (such as CLD [97] in various material lattices), values of B=s may range from 1032 m2 to 1036 m2 [36,37, and unpublished data from Dalton and coworkers]. The dominant decay mechanism appears to involve singlet oxygen (although this remains to be proven in detailed mechanistic studies). Photostability improves with exclusion of oxygen, increased lattice hardness, and the use of chemical and physical quenchers of singlet oxygen. Also, steric protection of chromophore sites vulnerable to singlet oxygen attack also significantly improves photostability. Photostability is improved at telecommunication wavelengths relative to that at shorter wavelengths consistent with the figure-of-merit definition. As with other properties of organic electro-optic materials, realization of adequate photochemical stability is possible with attention to detail. A final stability issue involves stability in the presence of high-energy radiation (gamma radiation and high energy (e.g., 25.6 MeV) protons). Several preliminary experiments have been carried out with little evidence of device performance degradation [45]. The resonance-stabilized p-electron system of chromophores may act to rapidly dissipate charge perturbation caused by the passage of ionizing radiation through electro-optic materials. However, such studies are difficult to execute and it will be difficult to assign any observed change in performance strictly to ionizing effects within the electro-optic material. Clearly, space qualification is an activity requiring further investigation.
6.8
Lattice Hardening
With the exception of polyimides and polyquinolines, few polymers exhibit sufficiently high glass transition temperatures to permit realization of adequate long-term stability of electro-optic activity for chromophore=polymer composites. A number of problems are encountered with such high temperature composite materials including sublimation of chromophores at high processing (e.g., poling) temperatures and poor solubility in traditional spin-casting solvents. While APC used extensively by Lockheed Martin researchers may provide both acceptable processability and long-term thermal stability, another approach is to increase the glass transition temperature of the final device material by crosslinking. This has the advantage of permitting good solubility and reasonable processing temperatures to be achieved in spin-casting and electric field poling. Various approaches to crosslinking have been reviewed at different times in the past 15 years [98,99]. A number of attempts have been made to employ photocrosslinking to achieve lattice hardening. This approach has the advantage of permitting the crosslinking event to be decoupled from temperature-dependent electric field poling. Unfortunately, competition for absorption of light by the electro-optic chromophores and photochemical initiators has, to this point in time, prevented adequate crosslinking densities from being realized. Thus, thermally activated crosslinking chemistries have largely been the focus of lattice hardening in more recent times. Two approaches have in particular become popular and these can be characterized as (1) reversible and (2) irreversible crosslinking reactions. The reversible Diels–Alder reaction can be used to good advantage by providing enhanced lattice hardening and by permitting good chromophore mobility at the glass transition (where crosslinks are disrupted). Choice of diene and dienophile, together with control of segmental flexibility, permits the glass transition temperature to be systematically varied with use of reversible Diels–Alder chemistry. This is very useful for temperature tuning in nanoimprint lithography [44] and in adjusting the glass transition temperature to the activation temperature of irreversible, thermally-activated crosslinking reactions. An example of irreversible, thermally activated crosslinking is illustrated in Figure 6.14 for the reaction of fluorovinyl ether to yield a cyclobutyl crosslink. Such chemistry can yield final glass transition temperatures of 2008C [27,100,101]. Moreover, high thermal stability can be obtained without attenuation of poling-efficiency or optical transparency provided appropriate attention is paid to processing conditions [27,100]. In particular, it is critical to adjust poling temperature and the temperature required for the thermal activation of the crosslinking to compatible values. Poling efficiency will be compromised if the temperature for the activation of crosslinking is lower than the
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Organic Electro-Optic Materials
0.5 O
O
N
F
F 0.5
O
O
F
O
F
F F
O
N
O
O O
O
PMI-TFV
TCBD-TFV S
F F
NC
O
O
NC NC
F
O CN
O
F
F
F
O
F
F
F
F
F F
O
F
F
F
F O
O F
F
F
O F cycloaddition
O
F F
F F
FIGURE 6.14 Irreversible crosslinking of the fluorovinyl ether functionality to yield cyclobutyl crosslinks is shown. In the example shown, crosslinking increases the glass transition temperature from 1508C to 2108C.
poling temperature. If the converse is true, insufficient lattice hardening (glass transition temperature elevation) will be achieved. Achieving sufficient lattice hardening is important for several reasons in addition to securing necessary thermal stability for the poling-induced electro-optic activity [27–31,100–102]. As noted in the previous section, photochemical stability and optical loss (associated with material processing) are frequently strongly dependent on lattice hardness.
6.9
Fabrication of Devices and Circuits: Electron Beam, Reactive Ion, and Photo-Etching
Waveguides have been fabricated by electron beam etching (EBE), RIE [7,83,84], photobleaching [98,103], two-photon photolithography, soft and nanoimprint lithography (as discussed in the next section), poling [104], and even by stress [105]. EBE is not financially practical for the production of commercial devices although it is useful for fabricating prototype devices particularly for prototype silicon photonic circuitry and devices. Photobleaching has been frequently used to fabricate devices such as Mach Zehnder modulators and to carry out postfabrication trimming [106–110]. Two-photon initiation of crosslinking reactions has been used to fabricate waveguide structures although such experiments to the present have been solely proof of concept rather than involving the demonstration of practical devices. The workhorse method, of fabricating practical waveguide structures in both organic and silicon materials, continues to be RIE. The crucial issue in fabrication is optical loss. For stripline structures, this has largely been a matter of controlling the kinetic energy of the reactive ions and of the hardness of the material being etched. For ring microresonator devices, mask resolution is an issue and surface roughness associated with mask resolution is frequently an important source of optical loss.
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6.10
Conjugated Polymers: Processing and Applications
Fabrication of Devices and Circuits: Soft and Nano-Imprint Lithography
Soft and nanoimprint lithography has been employed to fabricate stripline (Mach Zehnder), ring microresonator, and photonic bandgap devices [42–44]. Reversible Diels–Alder chemistry has been used to adjust glass transition temperatures to desired values. The performance of both stripline and ring microresonator devices fabricated by nanoimprint lithography was not noticeable degraded from that of the EBE master. Such processing potentially opens the way to low-cost, mass production of sophisticated circuitry.
6.11
Stripline, Ring Microresonator, Spatial Light Modulator, and Photonic Band Gap Devices
Electro-optic materials are typically incorporated into one of four basic device structures: (1) Stripline devices [7–9], (2) ring microresonators [25,39–42,109–115], (3) spatial light modulator (SLM) etalon, prism, or cascaded prism devices [116–118], and (4) PBG devices. Different device structures may be used to achieve different functions (e.g., optical switching, electrical-to-optical signal transduction, optical beam steering, radiofrequency–microwave beam steering, active wavelength division multiplexing (WDM), analog-to-digital conversion, reconfigurable optical add–drop multiplexing (ROADM), etc.) or to realize a special device feature such as compact size. For example, optical switching can be accomplished with any of the four basic device structures. The same optical=electrical field interaction lengths can be achieved in stripline and ring microresonator devices with devices of radically different sizes. For example, a stripline device might have a linear length of 0.5–3 cm while a ring microresonator device would likely have a diameter no greater than 500 mm (the size of a human hair). Ring microresonator devices thus permit a much more compact device size (e.g., a diameter as small as 80 mm) to be implemented. However, issues such as optical loss and bandwidth are also different for these different device structures. Ring microresonators are resonant devices (the optical beam constructively interferes with itself as it repeatedly transits the confining ring) and the bandwidth of the device is limited by the quality (Q) factor of the ring microresonator. Optical loss issues are different in these twodevice structures: With stripline devices, one is concerned with coupling and waveguide propagation losses while with ring microresonator devices, one has the additional concern of bending loss. Moreover, mask resolution in the fabrication of device structures will be an issue of greater concern for the fabrication of ring microresonator structures. Ring microresonators effect optical switching by acting as voltage-controlled optical bandpass filters or ROADMs. Only certain wavelengths of light (l ¼ neffL=m where neff is the effective index of refraction, L is the path length around the ring, and m is an integer) are coupled into the ring to form standing waves in a ring microresonator with the selection criteria defined by the effective ring circumference (neffL) and the index of refraction difference between the input and ring waveguides (the optical impedance matching condition). By application of a control voltage, a wavelength can either be passed or dropped. With stripline devices, switching is typically accomplished utilizing directional coupler devices [7–9,24]. An applied voltage modulates the coupling between two waveguides in close proximity, leading to voltage-controlled routing of optical signals between the two output ports. Because ring microresonators act as voltage-controlled optical filters, they can be used in a variety of ways including for active WDM, active optical interconnect reconfiguration, and laser wavelength tuning. As with any resonant device, the modulation bandwidth is limited by the photon lifetime in the ring or the optical bandwidth of the resonator. A device figure-of-merit for ring microresonators is the 3 dB bandwidth (frequency tuning) sensitivity factor (expressed in GHz=V). Alternatively, this can be expressed as a wavelength sensitivity factor (expressed in nm=V). This factor is related to material electro-optic activity, r33, by the following expression: BW( ¼ DnFWHM)=VFWHM ¼ Knn3r33=2neffd where VFWHM is the voltage required to shift the resonance an amount equal to the resonance width at half
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maximum, K is the confinement factor, n is the refractive index of the electro-optic material, neff is the effective refractive index of the ring, and d is the electrode spacing. Note that this relationship is derived from the index of refraction change of an organic electro-optic material upon applying an electric field in the direction of the aligned chromophores: Dneff ¼ K(n)3r33V=2d. Organic electro-optic materials currently yield tuning sensitivity factors on the order of 1–10 GHz=V but values of greater than 30 GHz=V should be possible with new organic electro-optic materials discussed elsewhere in this chapter. Optical loss is a major consideration in the design of ring microresonator devices. Bending loss will be an important consideration and will depend upon both the size of the ring and the index of refraction contrast between the core and the cladding of the ring microresonator device. This is illustrated in Figure 6.15 for parameters typically assessable with organic electro-optic core and passive cladding materials. The influence of index of refraction contrast on the free spectral range (FSR) is also shown in Figure 6.15. Material (absorption and scattering) loss will affect the Q (Unloaded Q ¼ 2.73105 neff=al where a is the sum of all loss (material, bending, and scattering due to wall roughness) in dB=cm and l is the wavelength of the resonant mode in microns; also Q ¼ n=D n where Dn is the width of the optical mode supported in the ring and n is its fundamental frequency or equivalently l=Dl where Dl is the width in nanometers and l is the central wavelength of the optical mode). Wall roughness (commonly defined by mask resolution) will also significantly impact the observed Q through its effect on a. Typical Q values realized with organic electro-optic materials lie in the range 105–106. This corresponds to Dl widths for the resonant modes of 0.2–0.02 nm or equivalently to Dn widths of 2–20 GHz. Q factors are very important for the operation of devices such as ROADMs since the critical device action is tuning the pass band center frequency by one or more Dl or Dn with a finite (small, e.g., 1 V) voltage. Clearly, a Q value approaching 106 is required for effective device operation with currently available materials but Q values of 105 would be more than adequate for new electro-optic materials with electro-optic coefficients of 300 pm=V. Other parameters of interest in understanding the performance of ring microresonators include the FSR, which is the distance between resonant modes, (FSR ¼ lmþ1 lm ¼ l2=ng(l)L where ng(l) is the group index given by ne(l) l(dne(l)=dl)), and the finesse (f ¼ 2p=aL). In addition to single ring microresonator structures, coupled ring microresonator structures have been fabricated [114]. These amplify the tuning sensitivity factor by the Vernier effect and amplification factors of 40 have been demonstrated. Coupled ring microresonator structures also lead to enhanced side mode suppression. Recently, a novel new architecture for multiple ring microresonators coupled to linear waveguides has been demonstrated [115]. 45
350
40
300 250
30 25
200
20
150
15
Radius (µm)
FSR (nm)
35
100
10 50
5 0 0
0.1
0.2
0.3
0.4
0.5
0 0.6
Index contrast (nco−nd)
FIGURE 6.15 The variation of free spectral range (FSR), the distance between resonant modes, and ring radius with the index of refraction difference between the core and the cladding is shown. As ring radius decreases, optical loss increases and acceptable ring radius sizes are normally chosen to represent the point where bending loss becomes comparable to other sources of optical loss.
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The relationship between device drive voltage and electro-optic tensor elements (r33 and r13) will vary with device structure. For example, for stripline Mach Zehnder interferometers, the drive voltage (or voltage required to produce a phase shift, f, of p) is given by Vp(MZ) ¼ lh=n3r33LG. For a stripline birefringent modulator, Vp (r33=r33r13)Vp(MZ). The drive voltage for a stripline directional coupling is a factor of 1.7 greater than that for a Mach Zehnder modulator [24]. Cascaded prism and superprism devices have been fabricated [8,9,116–118]; however, the electrooptic materials used in such devices required significant voltages to achieve wide angle beam steering. If current electro-optic materials (r33 300 pm=V) can be effectively incorporated into etalon and cascaded prism devices, operations such as optical switching can be accomplished with digital level drive voltages. A number of sophisticated device structures (phased array radar, A=D converters, spectrum analyzers, high bandwidth signal generators, etc.) have been demonstrated [7–9] but the discussion of these devices is beyond the scope of this chapter as is the discussion of modification of Mach–Zehnder modulators to improve linearity (spur-free dynamic range [SFDR]) and achieve wavelength-insensitive biasing. Indeed, ring microresonator structures have recently been used to significantly increase the SFDR of Mach–Zehnder modulators. While organic electro-optic devices have not received the concentrated engineering attention of lithium niobate electro-optic devices or of GaAs and InP electroabsorptive devices, an ever-increasing array of engineering advances are being demonstrated. These advances are best followed through reading the publications of IEEE, SPIE, AIP=APS, and OSA.
6.12
Conformal and Flexible Devices
Conformal and flexible devices can be conveniently fabricated by lift-off techniques [38]. As with OLEDs, the performance of organic electro-optic devices appears to be surprisingly immune to repeated high degrees of bending or flexing [38]. Indeed, no change in properties such as drive voltage, optical loss, etc., is observed until bending radii on the order of 1.5 mm are reached. The bending loss in flexed devices appears comparable to that in fabricated ring microresonators, so the act of flexing does not appear to cause any additional problems.
6.13
Integration with Silicon Photonics
Due to the large index of refraction of silicon (leading to a large index contrast between core and cladding), it is possible to fabricate ring microresonator and PBG structures of reduced dimensions relative to the all-organic device structures discussed earlier in this chapter. Ring microresonators with diameters of a few microns can be fabricated with acceptable bending loss [25]. Moreover, electrode spacings can be as small as 70–120 nm [25]. Indeed, with silicon photonic circuits, it has been demonstrated that it is possible to propagate light around a right angle (908) bend [119]. The small dimensions of silicon photonic circuitry together with the attractive prospect of integrating silicon electronics and photonics on a single chip exploiting CMOS manufacturing capabilities provide strong motivation for research and development. Indeed, researchers at Intel have demonstrated that electrooptic modulation can be effected in all-silicon photonics [120]. The problem with the demonstrated modulator is that bandwidth is limited and insertion loss is high. Here we discuss the integration of organic electro-optic materials with silicon photonic circuitry [25,41]. In the next section, we will discuss the integration of all-organic active photonic circuitry with silicon VLSI electronics [7]. Integrating organic electro-optic materials into silicon circuitry characterized by submicron dimensions would initially appear highly unlikely due to the poor enthalpic interaction between organics and pure silicon, and unfavorable entropy associated with nanometer scale structures. In reality, silicon photonic structures probably are best described as having a thin oxide layer coating on the surface and organic electro-optic materials thus interact strongly with such
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surfaces. In short, silicon photonic device structures can be filled with organic electro-optic materials by a variety of techniques ranging from spin coating to vapor deposition. A more serious problem is that of inducing acentric chromophore order necessary for macroscopic electro-optic activity. Silicon has a very low dielectric breakdown (e.g., 70 V=mm) and poling field strengths typically used with all organic device structures are unacceptably high. Doped silicon can be used as an electrode structure but exhibits poor conductivity; for example, such electrodes would not be acceptable for high-frequency modulation. On the positive side, the reduced dimensions of silicon photonic structures permit reduced electrode spacing, which translates to both reduced poling and drive (Vp) voltages. Effective electrode spacings, as small as 100 nm, have been achieved. BW=VFWHM sensitivity factors are currently about an order of magnitude greater in organic electro-optic=silicon photonic ring microresonator devices than in all-organic ring microresonator structures. Values on the order of 10 GHz=V have already been demonstrated and values greater than 60 GHz=V are likely with new electro-optic materials. Another challenge associated with organic electro–optic=silicon photonic hybrid devices is coupling significant optical intensity into the organic electro-optic materials from higher index of refraction silicon waveguides. There are two routes to the fabrication of hybrid device structures. The first approach makes use of evanescent coupling from silicon waveguides into an organic electro-optic cladding. The second approach places organic electro-optic materials into the slot of silicon photonic split ring microresonators. With silicon PBG devices, the vacancies in the silicon structure (including engineered defects) are filled with organic electro-optic materials. Theoretical analysis of electromagnetic wave propagation in split ring microresonators and PBG devices shows that a low loss transition of the optical mode occurs from silicon photonic waveguides to the organic electro-optic material concentrating high field intensity in the organic material. This high optical field intensity translates into effective electro-optic activity and to optical rectification [25]. In addition to electro-optic tuning, thermal tuning can also be utilized. One of the serious difficulties of silicon photonic circuitry is the issue of coupling light (e.g., from silica fibers or a laser source) into the silicon photonic circuitry. Waveguide grating (AWG) coupling appears to afford a viable solution to this problem. Other coupling schemes are also under investigation.
6.14
Integration with VLSI Electronics
Organic electro-optic waveguide devices have been horizontally and vertically integrated with VLSI semiconductor electronics [7–9,121–130]. Vertical integration offers a compact opto-chip structure but requires use of a planarizing polymer layer to reduce the large surface structural variations of a VLSI chip. To make connection from the VLSI electronics to the drive electrodes of the electro-optic waveguide device, deep vias (channels) must be etched through the planarizing polymer layer. This is conveniently accomplished using CF4 RIE together with spin-on glass technology. Vertical integration has been demonstrated and has been achieved without degradation in the performance of either electronic or optical circuitry.
6.15
Terahertz Signal Generation and Detection, Optical Rectification, and Electro-Optic Sensors
In addition to electro-optic activity, SHG and optical rectification can be observed for second-order nonlinear optical materials. With electro-optic activity, it is a low-frequency (typically, radiofrequency, microwave, or millimeter wave) field that causes the perturbation of the charge distribution (index of refraction) of the material that in turn influences light transiting the material. With SHG and optical rectification, it is the electric field of the optical beam that causes the nonlinear perturbation. The high optical fields existing in ring microresonators and other resonant devices can lead to optical rectification including high bandwidth rectification [25]. Terahertz generation and detection [12,13] is another
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manifestation of optical rectification arising when an intense pulsed laser beam (e.g., 45 fs, 30 mJ) is focused on a thin (70–130 mm) film of organic second-order nonlinear optical material [12]. In addition to poled polymeric thin film materials, organic crystals such as DAST and a variety of inorganic crystalline materials (ZnTe, GaAs, GaP) have been used for terahertz generation and detection. An advantage of poled polymer materials is the absence of phonon modes common to crystalline materials. Such modes limit operation over the 1–5 THz band. THz materials are useful for a variety of applications including imaging (biomedical, homeland security screening for plastic weapons, etc.) and spectroscopy. Sensors frequently rely on optical phenomena for detection and reporting. There are two types of optical sensors: (1) Those based on reporting changes in absorption or emission (e.g., fluorescence) and (2) those based on reporting changes in index of refraction. The most widely used technique based on exploiting index of refraction changes is surface plasmon resonance (SPR). Ring microresonators can also be used for sensing a variety of phenomena including strain [131], temperature, microwave radiation, and binding of chemicals (small molecules) and biomolecules. Use of electro-optic ring microresonators (or SPR) permits quick translation of the sensed phenomena into a proportional voltage through feedback control. Obviously, temperature can be sensed through the temperature dependence of the index of refraction and through the dependence of ring microresonator dimensions on temperature. Microwaves can be sensed either directly using an amplifying device such as a Luneberg lens or can be sensed via microwave-induced heating.
6.17
Applications and Commercialization
It is well appreciated that business models have dramatically changed over the past two decades as the full impact of globalization has come into play. Product development seldom occurs vertically within a given large corporation but horizontally involving the packaging of components from a number of individual vendors. Venture capital (start-up) corporations are playing an increasingly important role in the transition of new products to the marketplace. Moreover, nontechnical issues such as market demand can dominate commercial realities, e.g., telecommunications is still suffering from the ‘‘nuclear winter following the telecom bust’’ that dates from the beginning of the millennium. Thus, while telecommunications was viewed a major market for electro-optic modulators in the late 1990s, the need for increased telecommunications bandwidth was clearly overestimated and it will likely be several more years before the telecommunication market for electro-optic technology recovers. Defense and computer industries are likely to offer greater short-term prospects via utilization opto-chip technology for electrical–optical interconnection (e.g., backplane interconnection in the computer industry and sensor–processor interconnection in airborne and satellite platforms). There is the old joke that the ‘‘C’’ in CMOS stands for ‘‘cost’’. Very likely cost will be a dominant factor in defining the commercial application of organic electro-optic devices. As already noted, packaging and processing costs, rather than basic materials costs, will dominate final device costs. The unique processing options and ease of integration with disparate materials evident for organic electro-optic materials may be critical factors in controlling cost and influencing their utilization in established information technology (silicon CMOS) platforms. The potential for integration directly into silicon photonics is particularly intriguing in permitting a high density of devices on individual chips. The low power consumption (derivative from low drive voltage requirements) of organic electro-optic devices may be important for the practical utilization of such integrated structures. Another important application of organic electro-optic materials and devices will likely be that of niche applications that depend on some unique feature of organic electro-optic materials such as high bandwidth, low drive voltage, flexibility, light weight, etc. An example would be lightweight, flexible antenna (e.g., phased array) transmitter–receiver systems. The ease of fabricating special device structures, such as microresonator and photonic bandgap structures, from organic materials may be quite relevant. In like manner, the ability to fabricate three-dimensional photonic circuitry could impact special device applications.
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To this point in time, there is no evidence of significant commercial sales of organic electro-optic modulators. The one publicly traded start-up company offering such devices (see Figure 6.16) is Lumera Corporation in Bothell, Washington (NASDAQ, LRMA). Lumera appears to be a reliable source of both simple modulators and of electro-optic materials. Larger firms such as Lockheed Martin, Boeing, and Intel have viable research efforts and Lockheed Martin has continuously maintained a research effort in this field for two decades. As evident from the discussion in earlier sections, Corning has investigated the synthesis of variants of CLD- and FTC-type chromophores and of the photostability of such chromophores. Researchers at Boeing have demonstrated the viability of fabricating ROADMs employing organic electro-optic=silicon photonic split ring microresonator device structures. Of course, such research laboratory demonstrations are a long way from demonstrating a commercially viable system. Considerable material and processing refinements will be required before a commercially relevant system can be realized. Obviously, the potential for future commercialization may depend on material properties. Recent improvements in electro-optic activity to 300 pm=V, together with further anticipated improvements, should greatly increase the attractiveness of organic electro-optic materials. However, critical for commercialization will be the simultaneous realization of all additional auxiliary properties such as optical loss, thermal stability, photochemical stability, and processability. Again, good progress appears to be being made but the timescale of that progress will likely be critical to commercialization, which will also depend on developing auxiliary materials such as cladding and packaging materials. Indeed, development of an oxygen and moisture impermeable barrier (packaging) material appears important for the commercial success of both organic electro-optic and organic light-emitting devices. Again, the great advantage of organic materials is that their structures can be systematically varied to achieve the desired properties; however, such modification takes time and the real issue is whether or not materials with spectacular properties can be achieved on a timescale imposed by the marketplace. The focus of Lockheed Martin on the rational design of materials with low optical loss, of Corning on the design of material with adequate photostability, and of Lumera on the design of materials with adequate thermal and photostability are important steps in accelerating the transition of organic electro-optic materials and devices to commercial success.
FIGURE 6.16
A photograph of a Mach–Zehnder modulator manufactured by Lumera is shown.
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Acknowldegments The author wishes to thank his faculty colleagues (particularly Professors Bruce Robinson, Alex Jen, Philip Reid, Antao Chen, Axel Scherer, and William Steier), graduate students, and Drs. Susan Ermer and James Grote for many helpful discussions. The author also gratefully acknowledges the Air Force Office of Scientific Research and the National Science Foundation (DMR-0092380 and DMR-0120967) for partial support of this research.
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38. Song, H.C., M.C. Oh, S.W. Ahn, and W.H. Steier. 2003. Flexible low-voltage electro-optic polymer modulators. Appl Phys Lett 82:4432–4434. 39. Dalton, L., A. Jen, W. Steier, B. Robinson, S.H. Jang, O. Clot, H.C. Song, Y.H. Kuo, C. Zhang, P. Raibiei, S.W. Ahn, and M.C. Oh. 2004. Organic electro-optic materials: Some unique opportunities. Proc SPIE 5351:1–15. 40. Chen, A., L.R. Dalton, T.J. Sherwood, A.K. Jen. P. Rabiei, W.H. Steier, Y. Huang, G.T. Paloczi, J.K. Poon, A. Scherer, and A. Yariv. 2005. All-organic and organic-silicon photonic ring microresonators. Proc SPIE 5708:187–197. 41. Maune, B., R. Lawson, C. Gunn, A. Scherer, and L. Dalton. Electrically tunable ring resonators incorporating nematic liquid crystals as cladding layers. Appl Phys Lett 83:4689–4691. 42. Yariv, A., C. Zhang, L.R. Dalton, Y. Huang, and G.T. Paloczi. 2004. Fabrication and replication of polymer integrated optical devices using electron-beam lithography and soft lithography. J Phys Chem B 108:8006–8013. 43. Paloczi, G.T., Y. Huang, A. Yariv, J. Luo, and A. Jen. 2004. Replica-molded electro-optic polymer Mach-Zehnder modulator. Appl Phys Lett 85:1662–1664. 44. Firestone, K.A., P. Reid, R. Lawson, S.H. Jang, and L.R. Dalton. 2004. Advances in organic electrooptic materials and processing. Inorg Chem Acta 357:3957–3966. 45. Taylor, E.W., J.E. Nichter, F.D. Nash, F. Haas, A.A. Szep, R.J. Michalak, B.M. Flusche, P.R. Cook, T.A. McEwen, B.F. McKeon, P.M. Payson, G.A. Brost, A.R. Pirich, C. Castaneda, B. Tsap, and H.R. Fetterman. 2005. Radiation resistance of electro-optic polymer-based modulators. Appl Phys Lett 86:201122–1–3. 46. Liao, Y., B.E. Eichinger, K.A. Firestone, M. Haller, J. Luo, W. Kaminsky, J.B. Benedict, P.J. Reid, A.K.-Y. Jen, L.R. Dalton, and B.H. Robinson. Systematic study of the structure–property relationship of a series of ferrocenyl nonlinear optical chromophores. J Am Chem Soc 127:2758–2766. 47. Zhou, Y., L. Shaojun, and Y. Cheng. 2003. Poling properties of guest–host polymer films of X-type nonlinear optical chromophores. Synth Met 147:519–1520. 48. Kang, H., P. Zhu, A. Facchetti, and T.J. Marks. 2004. Self-assembled electrooptic thin films with remarkably blue-shifted optical absorption based on an X-shaped chromophore. J Am Chem Soc 126:15974–15975. 49. Sullivan, P.A., S. Bhattacharjee, B.E. Eichinger, K. Firestone, B.H. Robinson, and L.R. Dalton. 2004. Exploration of series type multifuctionalized nonlinear optical chromophore concept. Proc SPIE 5351:253–259. 50. Kang, H., A. Facchetti, H. Jiang, P. Zhu, and T.J. Marks. 2004. Synthesis and unprecedented electrooptic response properties of twisted p-system chromophores. Mater Res Soc Symp Proc. 833: 151– 156; Facchetti, A., G.R. Hutchison, S. Keinan, and M. Ratner. Control mechanisms for transport and nonlinear optical response in organic materials: a tale of twists and barriers. Inorg Chim Acta 357:3980–3990. 51. Professor B.H. Robinson of the University of Washington has initiated a systematic comparison of b values calculated by different theoretical methods with the assistance of theorists from around the United States. For simple variations of donors, acceptors, and bridges, various theoretical methods appear reasonably consistent in predicting trends; Dalton, L., B. Robinson, A. Jen, P. Reid, B. Eichinger, P. Sullivan, A. Akelaitis, D. Bale, M. Haller, J. Luo, S. Liu, Y. Liao, K. Firestone, N. Bhatambrekar, S. Bhattacharjee, J. Sinness, S. Hammond, N. Bruker, R. Snoeberger, M. Lingwood, H. Rommel, J. Amend, S. Jang, A. Chen, and W. Steier. 2005. Acentric lattice electro-optic materials by rational design. Proc SPIE 5912:A1–12. 52. Breitung, E.M., C.F. Shu, and R.J. McMahon. 2000. Thiazole and thiophene analogues of donor– acceptor stilbenes: Molecular hyperpolarizabilities and structure–property relationships. J Amer Chem Soc 122:1154–1160. 53. Casmier, D.M., P.A. Sullivan, O. Clot, K. Firestone, S. Lee, S. Heller, A. Brumbaugh, B. Millard, and L.R. Dalton. 2004. New paradigm in NLO chromophore design through a gradient bridge concept. Proc SPIE 5351:243–252.
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Introduction......................................................................... 7-1 Device Structures and Models............................................ 7-2 Characterizing Materials as Device Building Blocks . Device-Oriented Models . The Polymer Field Effect Transistor (FET) Structure . Contacts in OFETs . Extracting Mobility and Threshold Voltage for OFETs . Organic LEDs
7.3
Nir Tessler, Janos Veres, Oded Globerman, Noam Rappaport, Yevgeni Preezant, Yohai Roichman, Olga Solomesch, Shay Tal, Elena Gershman, Michal Adler, Vadim Zolotarev, Vladimir Gorelik, and Yoav Eichen
7.1
Polymeric Materials for OFETs ........................................ 7-11 Opportunities for Polymers in Organic Electronics . Material Challenges for OFETs . Transport Limitations and Strategies for Material Improvement . Polymeric Semiconductors . Polythiophenes . Stability and Extrinsic Factors . Deposition Techniques, Processing, and Film Morphology . Amorphous Semiconducting Polymers . Dielectrics and Interface Effects . Ambipolar Materials and Devices
7.4
Expanding the Frontiers of Organic Electronics Using Sequence Independent Synthesis Tools ................ 7-25 Nature’s Approach . Solid-Phase Synthesis of p-Conjugated Molecules . Electronic Peptides
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Summary and Outlook ..................................................... 7-37
Introduction
The field of organic optoelectronics has been constantly evolving since the late 1980s [1–6] picking up pace toward the late 1990s [7–13]. The evolution in this field has been largely driven and directed by the evolution and synthesis of new compounds [14–17] that enabled better devices [18–20] and sometimes even new functionalities [21]. From the device-engineering viewpoint, for a long time, there has been an advantage in using small molecules over polymers. This was due to the availability of small molecule building blocks as well as the better understanding of the characteristics of these blocks [22–24], compared to conjugated polymers. In recent years, this situation has started to change and in this contribution we try to address polymeric materials, in particular the relation between device engineering and material synthesis or development. We do that by examining how device engineering and analysis can support synthetic effort as well as discussing novel synthetic tools (and material characterization), which are specifically targeted toward improved device functionality. The functionalities we chose are (polymer) field effect transistors [25–28] and (polymer) light emitting diodes. 7-1
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This contribution is organized as follows: in Section 7.2, we mainly try to raise the reader’s awareness to the complexity of the physical picture underlying organic devices. We show that trying to find elegant physics, which applies across different device platforms and materials, is often a misleading notion and one also needs a strong engineering approach to make technological progress. In the following section (Section 7.3), we focus our attention on field effect transistors (FETs), which offer a promising platform for flexible display back planes and many other new opportunities. This part illustrates key factors for choosing or developing materials and discusses how much device performance is affected by the way materials are applied and processed. The importance of combining certain materials in device configurations is also examined together with the exciting progress achieved in this area. New polymeric materials will no doubt incorporate learning from the physics and processing knowledge accumulated. Some novel synthetic techniques described in Section 7.4 can play an important role in this by providing new toolkits for the controlled synthesis of polymers. These chemistry tools bring synthesis close to device engineering by using smart block assembly. Various approaches to solid phase synthesis of conjugated polymers and oligomers are reviewed including the biomimetic approach that utilizes electronic peptides. In Section 7.5, the chapter ends with concluding remarks and outlook for opportunities in the development and use of polymers in organic devices.
7.2
Device Structures and Models
Both the light-emitting diode (LED) and the FET structures can be found in classic textbooks discussing semiconductor devices [29]. The device structures commonly used are those of the P–N (or P-i-N) diode for the LEDs and the lateral metal oxide semiconductor (MOS) FET. These structures are based on (currently standard) processes of multilayer deposition as well as selective N and P type doping (see Figure 7.1) and rely on the mobility of 1 cm2 v1 s1 or more. In the organic semiconductor field however, some of the building blocks of the silicon technology do not exist or at best only partially present. This naturally imposes new limits on device fabrication procedures and device structures (see Figure 7.1c). For example, the low mobility of light-emitting organic semiconductors (typically 105–103 cm2 v1 s1) forces the thickness of LEDs to be on the
(a)
(b) Vg Gate
VS
Drain
Source P
VD
P i
P N
N
(c) Source
L
Drain
π-conjugated Insulator Conductor Vg
FIGURE 7.1 Schematic description of the MOS-FET structure (a) and of the P-i-N diode structure (b). (c) The organic FET (OFET) top contact structure, using a conjugated organic material in the channel.
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order of 100 nm, so as to avoid the need for unreasonably high applied voltage. Due to the low refractive index of conjugated molecules («r 2–3), the 100 nm scale is not much more than half of the emission wavelength thus making the metal induced fluorescence quenching and interference important mechanisms that must be closely controlled. The need to position the emission (recombination) zone within a fraction of a 100 nm imposes extremely high level of control over the electron and hole transport as well as over the injection properties of each. It was extremely difficult to produce such structures by solution coating without appropriate materials. As there were not enough building blocks (i.e., materials or processing) available and the ability to assemble (stack) blocks was limited, the inclination was to look for a magic material. This material has to be well matched to the available electrodes, have properly balanced electron and hole mobility, and exhibit high photoluminescence efficiency. By neglecting the idea of using building blocks we have actually left out the potential contribution of a device engineer and left it almost entirely on the shoulders of the synthetic chemist to develop a multifunction material. Small molecules make it possible to use building blocks, however, they are vacuum evaporated instead of the preferred route of solution processing or printing. Assuming that solution coatable building blocks are available, as becoming gradually more so recently, the next essential ingredient is a good understanding of the properties of these blocks so that one can design and engineer a structure by assembling functionalities. Defining and extracting these properties require some level of modeling that translates measured data (e.g., current or voltage) to a physical property (e.g., mobility, emission efficiency, or disorder parameter). We will try to illustrate below that although it is possible to establish such properties, one has to be very careful when extending the applicability of the physical models to device structures and materials.
7.2.1 Characterizing Materials as Device Building Blocks At a first glance, it seems that extracting the relevant device properties is very easy and all one needs to do is fabricate an LED and FET and test them. Indeed, many investigations are based on the idea that (a) by fabricating a thin device, as LED, one can directly measure the electric field dependence of the mobility; (b) by fabricating a long device, as FET, one can extract the charge density dependence of the mobility; and (c) by varying the temperature one can extract contact injection properties. Unfortunately, due to the mobility values and the device dimensions involved it is very difficult to separate the effects in a functioning device. As the material–device properties depend on its processing procedure and much thicker films require different processing procedure, one cannot extend the device dimension (as LED thickness) to the point where all these processes are strictly separable. To make real progress, a dedicated device-oriented model [30] is needed that would provide insight into the complex (nonlinear) problem at hand. The inevitable side is that as the modeling must account for the device structure and operating conditions, the extracted parameters are now strictly valid only for similar device structures and similar operating conditions (though as general guidelines they are still very useful). It is often useful in such models if some part of the device can be treated as an ideal component. For example, if it can be assumed that a contact is ohmic, or a material is trap free. Unfortunately, in practice it is rare that such assumption can reliably be made. A useful approach is to build test devices specifically designed to extract a certain property, however, one should make sure that the relevant device operating conditions are not altered (not an easy task). The other approach would be to perform a wide range of tests on a specific device so as to isolate the properties of interest. In organic field-effect transistors (OFETs), the use of high work function electrodes (for p-type OFET) and ideal dielectrics can help to reduce the number of variables.
7.2.2 Device-Oriented Models Most device-oriented models for organics share the assumptions found in common semiconductor device models [31]. The most obvious assumptions are (a) the carrier population is described as being at equilibrium (or quasi equilibrium); (b) the transport phenomena can be described using mobility (m)
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and diffusion (D) coefficients, which are treated as macroscopic parameters. The second assumption is more formally expressed as the transport can be described using its first two moments only (D and m). What we have in mind is a medium that can be described by two effective (mean) values and the challenge is in producing (through modeling) the values for D and m within the framework of interest. It should be obvious that by doing so we ignore some intellectually beautiful physics and we are left with a simplified picture that is, hopefully, just good enough for our purposes. There are many approaches to the modeling of transport and to the extraction of the mobility and diffusion coefficients from such models. The Monte Carlo simulation approach has proven to be very useful in understanding the relation between the mobility the electric field and the disorder parameter [32–35], in 707 TOF type of experimental configurations. The approach of averaging the transport through the most probable percolation path [36] proved to be useful in describing the charge density dependence of the mobility [37,38]. A Master equation approach was used to study the effect of both electric field and charge density on the mobility in a molecular semiconductor [39]. Recently, semianalytical models [40–42] helping to shed light on the complicated transport problem have also been developed. The semianalytical model by Roichman et al. [42] is based on the mean medium approximation (MMA) and is probably the only one that shares the assumptions found in standard semiconductor models (for better or worse). A percolation-type model, which predicts that there are bottlenecks or small regions in which the current density is much larger than average, do not agree with the assumptions behind the semiconductor device model equations. The assumptions behind the MMA model include 1. The semiconducting media can be treated as a uniform semiconductor having a density of states (DOS), which is defined by the distribution of the energies associated with the hopping sites, as would be found in a very large sample. 2. The charge distribution in energy follows a quasi-equilibrium distribution and can be described as the DOS multiplied by a Fermi–Dirac function. This model, adapted to organic materials, might be the only semianalytic approach that treats both charge density dependence and electric field dependence under the same framework (see Figure 7.2) and hence demonstrates the interplay between the two effects [30]. However, for the current discussion, a more important feature we can safely state is that whenever this model fails so does the standard
Charge density (cm-3)
Mobility (a.u.)
10–2
1015
1016
1017
1018
1015
1019
1016
1017
1018
1019 10–1
(b)
(a)
5.5x105
σ = 4 kT 10–3 σ = 5 kT 10–4
10–2
σ = 7 kT
10–5 10–6
E ≤ 1x105
Mobility (a.u.)
1014
σ = 4 kT 10–5
10–4
10–3
10–2
10–1
10–5
10–4
10–3
10–2
10–1
Relative charge density
FIGURE 7.2 (a) Calculated charge density dependence of the mobility for three different disorder parameters of a Gaussian DOS (electric field assumed to be negligible). (b) The effect of applied electric field (V=cm) on the density dependence of the mobility.
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semiconductor equation approach. This means that we can use a model that is adapted to organic semiconductors to check where the standard semiconductor equation approach must be used with some caution. Every time a device characteristic equation given as IDS ¼
W V2 mCins (VGS VT )VDS DS L 2
for FETs or 9 V2 JSCL ¼ «m 3 8 d for LEDs is used, the assumptions behind the semiconductor equations are invoked. The semiconductor equations [31] describe well the equilibrium transport in micron-thick films. However, it is not clear if the validity extends to organic thin-film (as OLED) devices. To answer this question we examine the organic material–adapted MMA model. Surprisingly, the MMA model’s assumptions are not inline with the real physical picture found in organic LEDs. The thickness of the active layer is of the order of 100 nm, making the volume being sampled by a carrier traversing the LED to be about (100 nm 30 nm 30 nm). This volume is very small and contains too few sites to reproduce the (disordered) DOS function that is found in a large sample (i.e., assumption 1 breaks). In addition, if we ignore the fact that the transport properties across the LED plane are nonuniform, the MMA model becomes less and less accurate as the applied voltage rises above 5 V, due to the resulting high electric fields. It has been shown [39], through the Master equation approach, that an electric field of 106 V=cm significantly distorts the charge density distribution in energy and it no longer follows the Fermi–Dirac distribution function. The situation with FETs seems to be more favorable because of the relatively low electric fields and the micron-scale distance between the source and the drain contacts. However, the interface phenomena that may be pronounced, depending on the processing procedure, in the FET channel are not accounted for. We should emphasize that although the standard semiconductor equations’ approach may not be strictly valid, it is, and probably will always be, very useful in capturing important features of LEDs, FETs, or photocells. Also, on this level, useful knowledge can be transferred from one device structure to the other allowing for accelerated development. However, one should acknowledge the limitations of these equations for in-depth understanding and fine optimization of real devices.
7.2.3 The Polymer Field Effect Transistor (FET) Structure When designing organic FETs, one has to make sure that no parasitic effects come into play. This is not only important for a device in a real application but also for transistors that are made only to extract p-conjugated material properties, such as mobility. For example, parasitic charging can easily be mistaken for redox-active materials that would exhibit hysteresis-like properties. Further in the section, we briefly mention issues related to the device layout, choice of contact materials, and the insulator material. Some of these issues can be found in textbooks [29,43] and are introduced or mentioned here just to make the picture more complete. The simplest OFET layout shown in Figure 7.1c can be described with the aid of the equivalent distributed circuit shown in Figure 7.3. The capacitors (and resistors) that are between the source and drain electrodes represent the channel region, whereas the outer elements represent the periphery of the device. As the gate electrode extends over the entire substrate, there can be significant parasitic current charging into the periphery.
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VD
RD
VS
R2
Rn
V n+1
VG
C ins
VG
Rs
R1 V2
VG
V1
VG
VG
VG
FIGURE 7.3 The equivalent electronic distributed-circuit of the CIp–FET layout shown in Figure 7.1c. The dashed capacitors at the sides represent the leakage to the periphery.
To avoid this charging effect, one can employ two textbook solutions [43], shown in Figure 7.4. In Figure 7.4a, the outer capacitors are eliminated by patterning the gate so that it will not extend beyond the region of the channel (and a bit under the contacts). Figure 7.4b shows the source and drain electrode layout designed such that the drain electrode will enclose the source electrode. In this case, the source electrode is not exposed to the periphery and the current measured through it is free of parasitic charging effects. A third solution, not shown, would rely on patterning the semiconductor as well.
7.2.4
Contacts in OFETs
The importance of the contact and the possible formation of a Schottky barrier [44] are becoming more important as the transistors are being made of higher mobility materials, which can support higher currents. As the current drawn from the contact is becoming larger, the voltage drop that is required to bias the Schottky diode also increases and the device performance is reduced. For example, Figure 7.5 shows the calculated potential distribution at the channel for a device structure similar to that of Figure 7.1c [45]. The device was biased in the linear regime (VG ¼ 5 V, VS ¼ 0 V, VD ¼ 3 V) and the potential and current distributions were calculated using a 2D device simulator. We note that for the 0.5 eV barrier height, the potential at the channel underneath the source shifts toward the gate potential by 0.7 V. As the transistor performance is dictated by the potential at the channel interface, this will also manifest itself as an apparent increase of the threshold voltage. This figure also demonstrates that in the top contact configuration, the loss in contact current-supply is compensated by the extension of the channel underneath the contact thus, drawing the current from a larger area [46]. On top of this effect, it has also been shown that the contact interface and its quality depend on the metal being evaporated in the organic layer [47] as well as the evaporation rate [48], and
(b)
(a)
Source
L
Drain
Π-conjugated Insulator
Vg
VS VD
FIGURE 7.4 (a) The patterned gate layout (b) Top view of the closed topology interdigitated source and drain contacts (note that the source in enclosed within the drain).
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Source
Potential at the Channel (V)
0 ∆F = 0.2eV
–0.5
VP_T= 0.7 V
–1 ∆F = 0.5eV
–1.5 –2 –2.5 –3 –3.5 –2
0
2 4 Length (µm)
6
8
FIGURE 7.5 Calculated steady state potential distribution at the channel interface. The applied bias is VG ¼ 5 V, VS ¼ 0 V, VD ¼ 3 V. The top and bottom lines were calculated for barrier heights of 0.2 eV and 0.5 eV, respectively. (From Tessler, N. and Roichman, Y., Appl. Phys. Lett., 79, 2987, 2001. With permission.)
other processing conditions [49,50]. It is not realistic to predict the effective barrier based on literature values of metal work functions as even inert metal surfaces are affected by contamination and processing history. Therefore, the barrier between the electrodes and the organic semiconductor energy level has to be examined carefully. As an example for a system that exhibits Schottky barrier effects we choose the electron conducting material C60 with the top contact electrodes of silver (Ag) and gold (Au). The energy level alignment based on literature values is shown in Figure 7.6. On the basis of this figure both the silver and gold should exhibit a Schottky barrier, with the barrier of gold being significantly higher. Figure 7.7 shows the source-drain current for two devices that were identical but for the top contact metal that was used. The structure was PþSi =50 nm SiO2=200 nm C60 = metal. Figure 7.7a shows the characteristics of the OFET with Ag contacts, which look ideal despite the high barrier predicted in Figure 7.6. Other measurements showed that the contact indeed imposes very little restriction on the current. The extracted threshold voltage was VT ¼ 1 V (see inset to Figure 7.7a) and the mobility was found to be m ¼ 0.5 cm2 V1 s1. Figure 7.7b shows same type of characteristics but for the OFET with Au-contacts. We note that the linear regime is replaced with a diode-like curve, which is a clear indication of a contact barrier. It is clear that the device is contact-limited and the effective threshold voltage was found to be VT_eff ¼ 2.4 V (see inset to Figure 7.7a). The effective mobility deduced in the deep saturation regime was only meff ¼ 0.06 cm2 V1 s1. Finally, we wish to point out that there is nothing in the OFET structure, shown in Figure 7.1c, that dictates whether the transistor would be electron or hole channel device. Typically, if there is a mobility imbalance between electrons and holes, it is much easier to produce a working device based on the higher mobility carrier. The choice of electrodes also dictates the carrier that is more favorable for injection into the channel. However, if the mobilities are not too unbalanced then the unfavorable injection (barrier) can be overcome, to some extent, by appropriate biasing (see Figure LUMO 7.7). In such a case, one can apply a bias 3.5 Ag 4.4 scheme where the source injects electrons C60 (VG>VS) and the drain injects holes 5.1 Au (VD>VG). This biasing scheme leads to light HOMO 6.2 emitting FETs in ambipolar devices, which we will discuss later. FIGURE 7.6 Energy level diagram for C60 silver and gold.
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0.15 VGS = 8 V
Current ^0.5
Current (mA)
0.1
0.05
0.02
(b)
0.015
Au
Ag
0.01 0
2
4
6
8
10
VGS
0.005
VGS = 4 V
VGS = 4 V 0 0
2
4 6 8 10 12 14 0 Drain source voltage (V)
2 4 6 8 10 12 Drain source voltage (V)
14
Current (mA)
VGS = 8 V
(a)
0
FIGURE 7.7 Source-drain current as a function of drain-source voltage for several gate-source bias conditions. The contact metals used were silver (a) and gold (b). The inset shows the transfer characteristics used to extract the effective threshold voltage.
7.2.5
Extracting Mobility and Threshold Voltage for OFETs
There are two important parameters used to benchmark OFETs, the threshold voltage and the carrier mobility. In disordered semiconductors (e.g., organics) there exists a charge density dependence of both the threshold voltage [51] and the mobility [37,42,52], thus making their extraction procedure [53–55] somewhat complicated. As the density of states may assume different forms (Gaussian, exponential, tail stated or their combinations) it is hard to predict a priori the functional dependence of the mobility on the charge density. This can be overcome by assuming a polynomial expansion relation [55] of the form m(VGS VT ) ¼
N X
Kn (VGS VT )n ,
n¼0
which for a strictly Gaussian DOS or a strictly exponential DOS reduces to a single power law [30,37]: m ¼ K0 [VGS VT]a. However, a technique for measuring the DOS in FETs was recently reported and indicated that the true DOS is a combination of a Gaussian and exponential density of states [28,51] (in semicrystalline materials, one would also need to consider mobility edge models) [56]. By developing the current equation using a density dependent mobility, one can devise a reliable procedure for extracting such a mobility [30,53,54]. If one chooses to use standard methods based on the transconductance of a density-independent mobility then the extracted values are generally wrong. Another approach for extracting the threshold voltage and mobility value was recently introduced [57] based on time resolving the source and drain currents. The idea behind this technique is based on the equivalent circuit shown in Figure 7.3 and simulations presented in Ref. [45] and this technique is shown in Figure 7.8. First the device is biased at, for example, VGS ¼ 0 V and VDS ¼ 3 V. At this bias, VGS and VGD are equal or above zero, such that injection of holes into the channel region does not occur at either contact. Next, the gate voltage is switched, at t ¼ 0, to VGS ¼ 2 V. Under these conditions, only the source contact has positive potential relative to the gate and hence charges flow from the source into the channel (the drain current is zero). As the channel becomes charged it gradually builds up toward the drain. Once the channel is almost completely full, some charges reach the drain and the current will flow out of the drain contact and soon after that the device reaches its steady state where the current that flows through the source exits through the drain (i.e., charging stops). If at t ¼ 0 the gate voltage is switched to VGS ¼ 5 V, then under these conditions both the source and the drain contacts have positive potential relative to the gate. These biasing conditions make the charge flow into the channel from both contacts and the channel builds up faster (see Figure 7.3
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VG t=0 IS ~ t –a
log (t ) IS ~ exp (–at ) IS = IDC
log (IS,ID)
t =0
FIGURE 7.8
tON
log (t)
Schematic description of the time resolved FET measurement.
in Ref. [45]). The point in time, where the charge populations injected from the source and drain contact meet is most pronounced in the upper curve of Figure 7.9 (marked by an arrow). As the actual voltage at which charges begin to be injected from the drain into the channel depends on VT (jVG VTj > jVDj), one can extract the threshold voltage from these measurements [57]. Figure 7.9 shows the measured source current as well as a fit based on simulating the transmission line equations.
7.2.6 Organic LEDs Organic LEDs have received significantly more attention than any other organic device structure. We do not attempt to present an overview of the large work conducted with respect to LEDs. Again, our main goal is to raise the reader’s awareness to the complexity of the physical picture and the issues faced when trying to apply models to describe materials and devices.
VG –2 V –3.7 V –5.3 V –6.8 V –7.5 V –8.3 V –9 V –10 V –12 V –13.6 V –15 V –16.6 V –18 V –20.8 V
Is (A)
10–6
10–7 10–4
10–3 t (s)
FIGURE 7.9 Experimental and transmission line simulation results (colored and dashed lines, respectively) of OFET charging current for different gate voltages (VD ¼ 8 [V]). All of the simulations are with the same parameters: VT ¼ 3.6 [V], m ¼ 1.4 104 [cm2=V s], and Inoise ¼ 9 108 [A]. (From Roichman, Y. and Tessler, N., Turn-on and Charge Build-Up Dynamics in Polymer Field Effect Transistors, San Francisco: MRS, 2005. With permission.)
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V
FIGURE 7.10 Schematic description of a thin film device being composed of a distribution of transport pathways (mobility values).
When using thin film devices one is often concerned with the possible occurrence of nonequilibrium conditions that broaden transient response making the mobility difficult to define [58]. It has been shown that due to the disordered nature of organic materials, TOF measurements performed on thin samples (below 1000 nm) typically result in dispersive current traces. It has been independently argued that a dispersive nature is because of the energy relaxation that takes place during the charge transport and depending on the degree of disorder can take 3–10 mm [32,58]. In such a case of continuous relaxation the definition of the mobility is problematic [58]. Also, this parameter (m) has been critically important to analyze device behavior throughout the field in a large number of publications. In the following text we suggest a physical picture that would accommodate both experimental observations. To gain another insight into the physical problem we refer to our previous discussion of the semiconductor equations and the MMA model. We suggested in Section 7.2.2 that the disordered nature of organic semiconductors would inherently lead to nonuniform transport properties across a thin film device. This idea has led to the proposition [59] that in very thin devices (below 300 nm) the dispersive nature is not because of dynamic relaxation but rather due to inhomogeneous distribution of the transport properties across the thin film. This implies that the film is composed of many parallel pathways each having a time-independent (i.e., well defined) mobility value (see Figure 7.10). The strength of this approach is that on one hand the notion of mobility is preserved and on the other, the dispersive nature is being accounted for. Using this picture, an analytic expression relating the measured transient current (I) to the mobility distribution function (g) can be deduced:[59] 2 d2 I(t) K d ¼ 3g dt 2 t Vt Here, K is a time independent factor determined by the experimental conditions and d2=Vt ¼ m. Applying this equation to experimental data, an electron mobility distribution function has been derived (Figure 7.11). Figure 7.11 shows that most of the sample is constituted of areas that exhibit an electron mobility of about 1109cm2 V1 s1. It also indicates that there are few regions that exhibit higher mobility values as would be expected for a disordered, nonuniform film [28,60]. Although the above picture raises some questions with respect to modeling of LEDs, it could also provide an answer to inconsistencies often found [42] in quantitative studies. It is also possible that the existence of fast and slow paths contributes to device degradation and, understanding these dispersive phenomena could help in screening for long-lived materials.
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Distribution function
7.3
7-11
Polymeric Materials for OFETs
7.3.1 Opportunities for Polymers in Organic Electronics
In Section 7.2, we tried to show how much caution one should exercise when applying existing device models for organics because of the specific physics operating in disordered thin films. The treatment becomes even more difficult when dealing with real polymers having the complexity of polycrystalline morphology, impur0 3x10–9 1.5x10–9 ities, and a variety of surface interactions. In the followMobility (cm2 V–1 s–1) ing text we examine these effects in practical materials and also look at the potential for the use of polymers in FIGURE 7.11 The electron mobility spatialdistribution function extracted analytically devices. from transient data of 300 nm thick MEHPolymeric semiconductors are exciting for the field of PPV based thin film device. (From Rappaport, electronics as they enable an entire new range of fabriN., Solomesch, O., and Tessler, N., J. Appl. cation techniques. They can be deposited from a soluPhys., 99, 064507, 2006. With permission.) tion: spin-, web-, or dip-coated, printed by inkjet, gravure or screen printing, to mention only a few possibilities [62–64]. The driving idea has been to pattern circuits by additive processes on large areas. This in turn opens up a range of new devices or potentially simplifies the fabrication of existing products. Today, polymeric semiconductors are investigated for a variety of applications including organic field-effect transistors (OFETs) organic or polymeric light emitting diodes (OLED, PLED), and organic photovoltaics (OPV). As was discussed in Section 7.1, different device structures rely on different material properties. For OPV cells, light absorption is essential and is accompanied by the dissociation of excitons into free carriers. In OLEDs, recombination of carriers via luminescence is required. It is often neither possible or nor practical to combine all of these requirements in one polymer even for a single application. In such cases, separate layers are devised with different functionalities employing different materials. For each of these layers, one has to design a new material, or building block, bearing in mind the electronic and physical interactions between them, including energy levels, optical and transport properties, and solubility differences. This task is complex and requires an interdisciplinary approach involving chemistry, physics, and analytical and printing expertise. The use of new fabrication techniques is not the only aspect making organic technologies exciting. Perhaps, the greatest advantage of polymeric and organic materials is the infinite variability of the molecular building blocks used for their design and synthesis. The freedom the technologist enjoys is much greater than the optimization of typical crystalline inorganic semiconductors via doping or alloying. It is possible, for example, to build copolymers with separate emissive and transport functionalities. One can affect the solubility or viscosity of polymers by a slight change of their substitution pattern. There is an increasing amount of knowledge accumulating on the electronic activity of newer and newer molecular groups, which has been helping to design better polymeric materials. During this development work, much attention has been paid to combining functionality with processability. In other words, more and more building blocks are being introduced, allowing for the engineering of better and more sophisticated device structures. It would be outside of the scope of this review to address all material functionalities for every type of application. The focus is limited to FETs and transport effects in polymeric OFET. Carrier transport is a key property that links materials for all applications. The learning in OLED and OPV devices has been extremely useful in providing guidelines for OFET materials and vice versa. In OLEDs, carriers are injected from the anode and cathode and they move through the polymeric film. Ideally, the rate at which electrons and holes are supplied, i.e., their mobility, should be similar. In OPV cells, carriers are separated and transported to the respective electrodes to create a photocurrent. In OFETs, carriers
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induced by the gate field move under the influence of the source-drain field. In each case, carrier transport is important and it is the driving semiconductor property, especially for FETs. OFETs have also become a workhorse to study transport in general in organic materials, although it will become clear that the interpretation of field effect data is not always straightforward.
7.3.2
Material Challenges for OFETs
The holy grail of OFETs has been to achieve mobilities similar to that of amorphous silicon by solution processing. Mobilities of around 1–5 cm2 V1 s1 [65] have already been demonstrated with small molecules such as pentacene, however, these require evaporation. Thus, there has been great interest to explore what is achievable with soluble, polymeric semiconductors. Recently, some polymers reached mobilities as high as 0.6 cm2 V1 s1[66]. These laboratory demonstrations, however, are often achieved in small number of devices in a controlled environment. For example, devices are fabricated and measured in N2 atmosphere to avoid the degradation of materials. Unfortunately, the highest mobility materials appear to be most sensitive to environmental effects such as oxidation and doping. The issue of stability needs to be considered for organic single crystal (OSC) materials in their solid as well as solution form for their practical use. Stability is particularly problematic for n-type materials [67] with high electron affinity. Chemical and electronic interactions with other layers also need to be considered. Ideally, stability should be achieved without the need for encapsulation, otherwise the cost and complexity of the device manufacturing process is dramatically increased. There is much work ahead to improve reproducibility of solution-coated layers, especially as the control of morphology for polycrystalline polymers is not yet ideal. This will be a critical issue for large area applications such as display back planes. To improve the yield of devices, the solution coated or printed OSC has to be highly uniform. The solvents used in the process should be safe to use, preferably nonchlorinated. The purity of polymeric materials and solvent formulations is much more difficult to control than for small-molecule OSC. Impurities in the latter are easier to remove by crystallization. It will probably never be possible to provide polymeric semiconductors with the same purity, as it is common in Si wafers. In currently used synthetic approaches, some impurities up to several percent may be present and they may well be acceptable as long as they are inactive. On the other hand, the detection of harmful substances may not be possible by analytical techniques and indirect, electrical characterization will be required instead. One of the realities of solution-coated, ambient-processed devices is that water will always be present in the final polymer coating and the device has to cope with it without the deterioration of the electrical properties. Recent development work in the field of polymer OFETs is beginning to address these issues. For example, bias stress has been studied for a number of polymers, sometimes as a function of the environment. Device stability with time is also being studied by a number of groups, identifying, for example, gradual doping of polymers due to ambient storage. Our understanding of these effects is just beginning to mature. It is clear that OFETs have improved immensely in the past decade in terms of the initial characteristics of best devices under well-controlled conditions [68–70]. The most serious challenge now is to translate the laboratory learning into better, practical materials that withstand the scrutiny of a real manufacturing process, yielding improved stability, reproducibility, and reliability. In the following text, the most important material classes and issues will be reviewed, not just in terms of mobility and electrical characteristics but also the demanding requirements relating to practical use.
7.3.3
Transport Limitations and Strategies for Material Improvement
It is well known that the carrier mobility is limited in organic solids. In a wide range of molecular crystals, the mobility appears to be limited to around 1–10 cm2 V1 s1 [71]. Recently, single-crystal rubrene OFETs were reported with the mobility of 15 cm2 V1 s1 [72]. The reason for the mobility limitation is that molecular materials are not covalently bound and electronic orbital overlap is limited. A robust organic material with the mobility of 1 cm2 V1 s1 would still be an interesting competitor to amorphous silicon (a-Si). Some examples of molecular semiconductors are shown in Figure 7.12.
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7-13
Pentacene Rubrene
R
S
S S
n
S
R
n = 1–3 R = alkyl (C6H13)
dialkyl oligothiophene
FIGURE 7.12
Small molecule and oligomeric semiconductors.
Vapor-deposited polycrystalline pentacene has become a widely studied small molecule semiconductor approaching the molecular crystal mobility limit with mobilities as high as 5 cm2 V1 s1 [65]. These rigid molecules readily crystallize in a uniform manner with an efficient pp stacking. Under ideal conditions, the molecules lie perpendicular to the substrate and the effect of grain boundaries is not significant in the first layers deposited on the substrate—which is at once the gate dielectric. Several other rigid rod molecules afford efficient crystal structures when vacuum evaporated. Experiments on oligophenylenes [73,74] and oligothiophenes [75] showed that the mobility may be enhanced by using longer oligomers, suggesting that longer conjugation length is an advantage. The six- and eightmembered 6T and 8T thiophene compounds resulted in mobilities above 0.1 cm2 V1 s1 [76]. Another strategy has been to create closer packing crystal structures, for example a variety of substitution patterns, which have been investigated [77]. In practice, one of the most effective methods for increasing crystallinity has been to introduce alkyl-end substitution. True plastic electronics would require an ambient processable semiconductor. Solubilizing strategies have been applied to both small molecules and rigid oligomers. To make pentacene printable, a variety of precursor molecules were tried [78]. Unfortunately, the chemical conversion process following solution coating is difficult to complete and ultimately, the mobility and device stability suffer. Soluble pentacene derivatives and heterocyclic oligomers have also been applied but they are still inferior in processability to polymers [79,80]. Solution-processed oligothiophenes demonstrate 0.01–0.1 cm2 V1 s1, however, the film morphology is extremely sensitive to the deposition conditions [81,82]. Most printing or solution coating approaches therefore have been focussing on polymeric semiconductors. Polymers offer improved uniformity and film forming. Polymeric versions of the above-mentioned oligomers also aim at increased conjugation for increased intrachain transport while maintaining ordering, i.e., crystallinity for interchain charge transfer.
7.3.4 Polymeric Semiconductors Polymeric OSCs of today have not exhibited mobility as high as the model small-molecule pentacene. In polymers, electronic coupling may be relatively strong within a chain, however, only in one dimension and for a limited distance. At the slightest twist of the polymer, delocalization is disrupted. Carriers face inevitably interchain transfer, which is limited by the poor overlap between neighboring chains. Ordered, crystalline domains are helpful to increase overlap, indeed polycrystalline polymers such as poly(3-hexylthiophene), P3HT, are used in OFETs. There has been a drive to increase mobility in polymers on both accounts, i.e., by increasing the conjugation within a chain and by making them crystalline. The early classic conjugated polymer, polyacetylene, is a good example for this, having alternating single and double bonds and also being crystalline. One of the first studies of field effect in
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polymeric materials relate to polyacetylene [83] (see Figure 7.13). Despite exhibiting high metallic conductivity upon doping, as a C C C C C C σ semiconductor, this polymer showed very low mobility (104 cm2 V1 s1) in subsequent studies [4,84]. As this polymer is π insoluble, the Durham route precursor approach was used for its C C C deposition, which results in high purity but low degree of orC C C dering. Interestingly, by introducing mesophase forming chain FIGURE 7.13 Polyacetylene, the ends, the ordering and thus the mobility could be increased to classic conjugated polymer. 3103 cm2 V1 s1 [85]. Polyacetylene is a good illustration that despite the long, alternating conjugated bonds, transport is hindered by twists in the chain. The band gap of this polymer is very much affected by the chosen precursor route for its preparation and chain ordering. Polyacetylene in an ideal crystalline structure would have a very small band gap. This material is also very sensitive to photooxidation and ambient doping, making it unfavorable for device applications. Practical polymers for OTFT devices require low levels of intrinsic or extrinsic carriers. In fact, most polymers discussed here in their clean state would be intrinsic semiconductors, with practically no free carriers in them. This is achieved by a reasonably large band gap and high material purity. Some examples are shown in Figure 7.14.
7.3.5
Polythiophenes
The effect of increasing molecular weight on transport has been well demonstrated on oligo- and polythiophenes [86]. The mobility of thiophene oligomers was reported to increase from terthiophene (3T) to octithiophene (8T) [87]. Longer thiophene polymers were made soluble by introducing alkyl pendant chains such as in poly(3-hexylthiophene) (P3HT). Regiorandom thiophene polymers, however, resulted in modest mobilities of only 104 cm2 V1 s1 because of the disrupted conjugation at head-tohead couplings [88]. In addition, the longer the solubilizing alkyl chains were, the lower the mobility appeared in an essentially disordered material [89]. With the introduction of the fully conjugated regioregular poly(3-hexylthiophene) (rr-P3HT) mobilities in the order of 0.1 cm2 V1 s1 were achieved [90,91]. The regioregular structure has a lower band gap. The polymer is also highly crystalline and exhibits lamellar structure, the thiophene rings being perpendicular to the substrate in its most effective configuration [19] (see Figure 7.15). To induce such a state, the substrate is typically treated with a hydrophobic self-assembled monolayer (SAM) such as hexamethyildisilazane (HMDS) or octatrichlorosilane (OTS). In addition to increasing the solubility, the alkyl pendant groups on P3HT and other polyalkyl thiophenes (PAT) facilitate the stacking on the substrate. Bulkier side-chains have been shown to inhibit ordering, as indicated by electron diffraction and low carrier mobilities in devices [92]. The conjugated rr-P3HT exhibits higher mobility when the molecular weight of the material is increased [93] (see Figure 7.16), a phenomena found also in less ordered polymers [60]. Interestingly, despite their lower mobility, the crystallinity of low molecular weight polymers was more pronounced (Figure 7.17). This was attributed to the greater influence of grain boundaries in the crystalline, low MW polymer.
7.3.6
Stability and Extrinsic Factors
Highly conjugated, electron-rich materials are sensitive to air doping, reducing on and off ratios, and moving the threshold voltage to positive values for p-type OTFT. P3HT, for example, requires inert atmosphere during both processing and operation for a high on and off ratio [94]. The doping mechanism is believed to be complex formation with oxygen [95]. The formation of the charge transfer complex can be reversed by heating devices in vacuum to remove oxygen, however, on returning to air the off-current rises drastically [96]. One approach to reduce this effect is to lower the highest occupied
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C6H13
C6H13 S S
*
S
S
n
S C6H13
*
*
C6H13
*
S
S
C10H21
C10H21
P3HT
C8H17
n
S
BTT
C12H25
C8H17
S S
*
S
*
n
S
S
n
S
* PQT12
*
C12H25
F8T2
C8H17
C8H17 N
S
*
N
C8H17 n
C8H17
* N
*
n
F8BT
*
TFB
*
n
*
N *
n
*
PF2/6 PTAA
FIGURE 7.14
Polymeric semiconductors.
molecular orbital (HOMO) level of the polymer by disrupting the conjugation either by twisting the polymer chain, or introducing other aromatic groups in the polymer backbone [97]. Examples of these polymers, poly(3,3 dialkylquarterthiophene) (PQT) and a thieno[2,3-b]thiophene containing polyquarterthiophene (BTT) are shown in Figure 7.2. The introduction of phenyl groups was also successfully used in oligothiophenes [62]. Unfortunately, low ionization energy also leads to easier photooxidation, which is a permanent effect [98]. In addition to oxidative effects, ionic impurities in the semiconductor, substrate, or insulator also severely affect the device performance, causing hysteresis [99,100]. The poor stability of some semiconducting polymers and their sensitivity to external factors often make it difficult to deduce true intrinsic properties of a certain material [101,102]. Most commonly, the off-currents in a TFT are governed by extrinsic factors and it also has an impact on subthreshold slope, threshold voltages, and bias stress. In fact, this sensitivity prompted some researchers to use OSC materials as gas and chemical sensors in TFT devices [103,104]. The observed selective sensitivity is related either to the chemical nature of the semiconductor functionalities or the interaction of the grain
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5 s s
s
a
a
s
s
s
b
s
s
b
Intensity (a.u.)
(b)
(a)
4.5 4 3.5 3 2.5
(010) (300)
1.5
(200)
1
(100) (010)
2
0.5 (100)
FIGURE 7.15 Two different orientations of ordered P3HT domains with respect to the FET substrate by wide-angle x-ray scattering. (From Sirringhaus, H., Brown, P.J., Friend, R.H., Nielsen, M.M., Bechgaard, K., Langeveldvoss, B.M.W., Spiering, A.J.H., Janssen, R.A.J., Meijer, E.W., Herwig, P., and Deleeuw, D.M., Nature 401, 685, 1999. With permission. Copyright Macmillan Publishers Ltd.)
boundaries with gases or vapors. Nevertheless, organic materials in TFT applications may stand a good chance to achieve acceptable lifetime when compared to light-emitting diodes or photovoltaic cells, where higher energy excited states emitting or absorbing light play a role.
7.3.7
Deposition Techniques, Processing, and Film Morphology
It is clear that the mobility and device properties using a particular semiconductor depend very much on the configuration of the device and other layer elements used in it such as the source and drain electrodes and the dielectric interface. The characteristics can also be very different for the same semiconductor depending on the deposition technique and processing, which in turn influence the microstructure. It is therefore critical to understand the film microstructure in a particular device. Typical tools used to evaluate film morphology are scanning electron microscopy (SEM), polarized optical microscopy, electron diffraction, x-ray diffraction, especially at low angles as the active interfaces investigated are extremely thin. Earlier it was mentioned that tendency for P3HT stacking depends on the surface treatment; hydrophobic surfaces driving the pendant alkyl chains toward the substrate. This process is also dependent on the solvent used for the deposition and the speed of drying [90]. High boiling point solvents such as cyclohexylbenzene (CHB) and 1,2,4-trichlorobenzene (TCB) have been shown to enhance mobility compared to chloroform [105]. (Table 7.1, Figure 7.18) Surprisingly, even the noncrystalline poly(2-methoxy-5-(30 ,70 -dimethoxy)-1,4-phenylenevinylene) (MDMO-PPV) or poly[2-methoxy-5-(20 -ethyl-hexiloxy)-p-phenylenevinylene] (MEH–PPV) were reported to be influenced by the solubility properties, which could be altered via the choice of solvent [106] or the molecular weight [60]. Coating from chlorobenzene resulted in a greater number of chromophore aggregates, facilitating interchain transport and higher mobility in TFTs. Increased currents in other emissive polymer layers and OLED structures have also been reported due to solvent-controlled aggregation [107]. The dynamics of film forming are critical for a polycrystalline semiconducting film. The speed of drying, the nature of the surface, and the way the solution is deposited all influence the microstructure. Drop casting, for example, resulted in stronger edge-on orientation and higher mobility for P3HT than spin coating [90]. This sensitivity to rate effects is very similar to that in vacuum deposited pentacene devices. In general, slower film formation coupled with elevated temperatures leads to higher degree of crystallinity, but it may also lead to the greater influence of grain boundaries. A recent study on a
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(a) Au source
Au drain
Gate voltage –80 V
P3HT
Drain current (mA)
SiO2
–0.2
Si Gate –70 V
–0.1
–60 V –50 V –40 V –30 V 0V
0.0 –40
0
–80
Drain voltage (V) (b)
Mobility (cm2/Vs)
10–2
10–3
10–4
10–5
Group A Group B Group C
10–6
2
4
6
8 10
20
40
Molecular weight (kD)
FIGURE 7.16 (a) Output characteristics of a rr-P3HT transistor (b) and the mobility as a function of average molecular weight (b). (From Bao, Z., A. Dodabalapur, and A.J. Lovinger, Synth Met., 18, 699–704, 1996. With permission. Copyright 2003 by John Wiley & Sons, Inc. )
different polymer, poly[2,7-(9,9-bis(2-ethylhexyl)fluorine] (PF2=6), demonstrated that different liquid crystalline phases [108] can be attained in the same polymer depending on the molecular weight [109]. Ong and coworkers [110] introduced a nanoparticle approach to coat a polymeric semiconductor. The authors dispersed PQT in dichlorobenzene by sonicating a hot solution and subsequently cooling it to room temperature. The idea was to predetermine the polymer ordering before coating it onto the substrate. Increased mobility was reported compared to hot solution coating. The microstructure may also be influenced by thermal annealing after deposition. Mobility in poly(9,9-dioctylfluorene-co-bithiophene) (F8T2) was shown to increase with annealing temperature [111]. The polymer film underwent a phase transition—at room temperature, the as-spun semiconductor films exhibit an isotropic, amorphous phase. On annealing at elevated temperatures, crystalline, and liquid crystalline phases were observed. The order in the material correlated with the field effect mobility. A liquid crystalline, oriented structure can be further promoted by rubbing the substrate [112]. An amorphous to crystalline transition was also demonstrated in a eutectic mixture of rubrene and a polymer. The mixture was amorphous when coated, but on heating and rapid cooling, large crystalline domains were grown and a mobility of 0.1 cm2 V1 s1 was achieved [113].
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100
100
s
s
s
s
s
s
s
Counts
010
S
200
010 S
300
Mn = 3.2 kD Mn = 33.8 kD
5
10
15
20
25
2q
FIGURE 7.17 X-ray diffraction (XRD) data showing a more intense <100> peak alkyl chain spacing for a lower MW P3HT film compared to the high MW polymer. The <010> peaks (p stack direction) were small and showed no evidence of preferential orientation. (From Kline, R.J., et al., Adv. Mater., 15, 1519, 2003. With permission.)
When the semiconductor [115,116] and the contacts [114] are printed (Figure 7.19), the morphology of the active layer is influenced by many more complex factors such as the temperature of solution and substrate, the print speed and drop-to-drop overlap in case of inkjet deposition. It has been demonstrated on a number of inkjet-printed OTFT devices that the process can be optimized to obtain at least as good results as by spin coating and even achieve decent uniformity of device performance [20]. The most widely used material for printing has been F8T2 [114,117]. Paul et al. printed F8T2 and a regioregular polythiophene, XPT (Figure 7.20), resulting in mobilities of 4103 and 0.1 cm2 V1 s1, respectively [116]. In most cases, printed devices greatly benefit from an increased on and off ratio due to the localized deposition of the OSC in the channel area. This reduces any stray currents, which are not controlled by the gate. How much the printing sequence may influence the device quality is well illustrated by techniques directionally depositing a semiconducting film. Coating polyalkylthiophenes by a friction transfer (drawing) method, the mobility becomes highly anisotropic, being higher in the direction of drawing the polymer [118]. It can be anticipated that a similar improvement can be achieved with most conjugated polymers.
TABLE 7.1
Bp (8C) Chloroform Thiophene Xylene CHB TCB
Device Parameters as a Function of the Solvent Used for Depositing P3HT HMDS Mobility (cm2=V s) 60.5–61.5 84 138–139 239–240 218–219
Ion=Ioff
Subthreshold Slope (V=dec)
FDTS Mobility (cm2=(V s)
0.012
105
2.45
0.042 0.022 0.12
103–104 106 106
5.5 1.8 1.7
Ion=Ioff 0.0076 0.030 0.041 0.049 0.063
Subthreshold Slope (V=dec) 104 104–105 105 105–106 105–106
9.1 6.5 5.1 4.25 3.8
Source: From Salleo, A., and M.L. Chabinyc. Stability of organic transistors. In Organic Electronics. Weinheim: Wiley-VCH Verlag. Copyright 2004, American Chemical Society. With permission.
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106
Intensity (a.u.)
105
s s s s
100
TCB, in-plane TCB, out-of-plane ch, in-plane ch, out-of-plane
104 010
200
300
103 102 101 0
0.5
1 q
1.5
2
(Å–1)
FIGURE 7.18 Grazing incidence x-ray diffraction measurements for P3HT thin films spin-coated from TCB and chloroform on HMDS treated substrates, with out-of-plane and in-plane scattering geometry. The (100), (200), and (300) plane reflections are due to the lamellar layer structure, and the (010) reflection is due to d–d interchain stacking. (From Salleo, A., and M.L. Chabinyc, Organic Electronics, Weinheim, Wiley-VCH Verlag, 2005. Copyright 2004, American Chemical Society. With permission.)
7.3.8 Amorphous Semiconducting Polymers The earliest application of polymeric semiconductor was in the 1970s, which was polyvinyl carbazole (PVK)used organic photoreceptors. In organic photoreceptors, it is preferred to employ an amorphous and optically clear film. PVK was also the first organic polymers in which electronic transport was studied in detail [119]. Blends of luminescent molecules with PVK have been the first solution-coated polymeric light emitting diodes [120]. In PVK, the active carbazole unit is a pendant group and due to the lack of conjugation in the main chain the mobility is low, only about 105 cm2 V1 s1. Amorphous polymers have relatively low mobility even when the backbone is conjugated (e.g., polvinyphenylenes, PPVs). Not surprisingly, their use have been mostly limited for OLEDs, where the low mobility is sufficient. One of the most widely studied class, polyfluorenes, also originate from OLED use [121]. Typical bulk mobilities in these polymers are in the range of 103 cm2 V1 s1, as measured by time of flight (TOF) [122].
Channel
Source/drain
Gate
200 µm
FIGURE 7.19 Image of ink-jet printed electrodes in a TFT of 5 mm channel length. The inset shows a whole image of the device. The distance between printed dots is changed from 15 to 50 mm to get target thicknesses (100–300 nm). (From Ong, B.S., et al., Adv. Mater., 17, 1141, 2005; Kawase, T., Shimoda, T., Newsome, C., Sirringhaus, H., and Friend, R.H., Thin Sol. Films, 438, 279, 2003. With permission.)
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(a) Print direction
400 µm (b) Prin
2
t di
rec
tion
Au contact
1
Printed polymer
µm 20 40 substrate
60
Au contact
FIGURE 7.20 (a) Micrograph of an array of OTFTs with printed polymer XPT on gold source and drain contacts defined using wax printing and etching. (b) Topography of a 0.45% solution of F8T2 printed on gold contacts (200nm thick) acquired by AFM in tapping mode. A drop size of 35-mm and overlap ratio of 50% was defined in the printing program. The line-to-line overlap of the printed polymer is clearly seen in the corrugation of the surface. (From Yasuda, T., K. Fujita, T. Tsutsui, Y.H. Geng, S.W. Calligan and S.H. Chen, Chem Mater, 17, 264–268, 2005. With permission. Revised from Kateri, E., P., William, S.W., Steven, E.R., and Robert, A.S., Appl. Phys. Lett., 83, 2070, 2003. With permission. Copyright 2003, American Institute of Physics.)
Amorphous layers of polyfluorenes and their copolymers have also been studied in TFT devices. F8T2 and TFB, for example, exhibited 4103 and 103 cm2 V1 s1 mobilities, respectively [111]. Some amorphous polymers, such as polytriarylamine (PTAA), however, have reached mobilities up to 102 cm2 V1 s1. Although this is probably close to the limit achievable with an amorphous polymer, they have a very useful role for OTFT development. Transport in amorphous polymers can be studied by
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a variety of techniques and this makes them a very interesting tool for understanding better organic transistors. PTAA polymers originated from the field of electrophotography, where their amorphous nature was of benefit. This material class was further developed for OFETs to capitalize on their isotropic character—excellent uniformity and stability [123]. An amorphous layer would yield the same mobility irrespective of the device configuration such as conventional or vertical TFTs, insensitive to a substrate template. It also yields the same mobility irrespective of the deposition technique, so that OFETs can be printed with high uniformity and yield. The same homogenous layer is formed from a variety of solvents, which makes it easier to formulate printing inks. The mobility of PTAA polymers increases progressively with molecular weight, up to 80–100 repeat units, beyond which the increase diminishes [124]. Interestingly, this occurs despite the effective conjugation length being only about 6–10 units. For longer oligomers, there is little change in the optical properties in solution or in film form. It was speculated that in a film, a higher molecular weight material may have statistically larger number of conjugated units not affected by twists in the chain. Recently, mobilities of 0.03 cm2 V1 s1 were reported for a PTAA copolymer, a record number for any amorphous polymer [124]. Studies on amorphous PPV derivatives gave further insight into field effect in amorphous polymers. It was found that the mobility was dependent on the gate voltage, i.e., the induced carrier density (see Figure 7.21) [60]. The mobility increase was attributed to the filling of lower energy states in the DOS [30]. For systems with three-dimensional transport and relatively low mobilities typically a strong increase was observed. The effect has been analyzed theoretically confirming that at higher charge densities the mobility increases. A number of device models were developed treating this problem (see earlier in this chapter). The increase depends on the shape and width of the DOS [28]. For example PTAA and TFB, which are likely to have a narrow Gaussian DOS, there is little dependence on the charge density [123,125]. Note that in some OFET devices, the mobility is reported to decrease with increasing charge density, which is not yet understood. One explanation of this mobility may be related to the effect of increased gate field, forcing the carriers to move through less favorable interface states changing the transport regime toward a two-dimensional case. Another explanation may be related to the shift of the threshold voltage [51] that would make the apparent mobility smaller. Charge density effects are probably the easiest to investigate on amorphous OSC, enabling a generally better understanding of transport modes operating in different device types and comparing bulk and
Charge Density (cm–2)
Charge mobility (cm2v–1s–1)
1011
1012
10–5
VDS = –1
10–6 0.1
1 Insulator potential drop (V)
10
FIGURE 7.21 Experimentally extracted hole mobility in an MEH-PPV-based FET. (From Tessler, N. and Roichman, Y., Org. Electron., 6, 200, 2005. With permission; Shaked, S., Tal, S., Roichman, Y., Razin, A., Xiao, S., Eichen, Y., and Tessler, N., Adv. Mater., 15, 913, 2003. With permission.)
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TFT mobility. In particular, the use of amorphous OSC made it possible to shine light on the influence of dielectric interfaces independent of morphology effects.
7.3.9
Dielectrics and Interface Effects
The performance of semiconducting polymers in an OTFT is very much dependent on the other materials used in the device. In particular, the interface between the gate dielectric and the semiconductor has been shown to be critical. The variety of dielectric materials employed in OTFTs was recently reviewed by Fachetti et al. [126]. As the active layer in a TFT is only a few nanometres, the microstructure of the interface region has a strong influence on transport. First of all, the roughness of the interface has to be minimal, which can be ensured via the choice of orthogonal solvents and avoiding intermixing between the dielectric and the semiconductor layers [69]. Valleys and peaks in the channel can act as traps for carriers and the gate field can significantly reduce escape probability perpendicular to the direction of transport. Rough surfaces also diminish the growth of large crystalline domains. Inorganic dielectrics have been widely used for bottom gate devices, as they are not affected by the choice of solvents. SiO2 has been the workhorse to study early OTFT devices. A key driver for dielectrics has been to increase capacitance and thus reduce the operating voltage. This can be done effectively by high-k inorganic dielectrics as was shown with a variety of materials such as TiO2, barium zirconate titanate (BZT), and barium strontium titanate (BST) [127]. These materials can be deposited either by sputtering or anodization. High-k inorganics have been effective for crystalline small molecules such as pentacene, but less successful for polymeric materials (for a detailed summary of high-k insulator results see Ref. [126]). In bottom gate devices, the dielectric also serves as the substrate. Unfortunately, the chemical groups found on surfaces of inorganics are often extremely variable and unpredictable. The wide range of chemical treatments, SAMs, used in the literature illustrate that the chemical nature of the interface has a strong impact. The local chemistry influences the morphology of the polycrystalline layer and also provides an electronically more uniform interface with transport sites of semiconductors. Possible effects due to SAM treatments have been reviewed in Ref. [25]. SAM chemistries with an electron accepting functionality can modulate the charge density in the channel and change threshold voltages by as much as 50 V [128]. A high capacitance dielectric with an ideal interface exhibiting uniform surface chemistry would be vital to study in further detail charge carrier density effects in OTFTs. Organic polymeric dielectrics offer a more practical way to build OTFT devices by solution coating. Examples of polar and nonpolar dielectric polymers are shown in Figure 7.22. Some early studies reported higher mobilities using high-k, polar resins such as cyano-ethylpullulan (CYMM) or poly(vinyl alcohol) (PVA) and attributed it to the increased capacitance [129,130]. Such relatively high-k polymers have been less favored recently. Unfortunately, high-polarity organic insulators are difficult to use effectively in devices without causing hysteresis by slow polarization of high-dipole molecular segments [25]. Often, such slow polarization leads to a larger capacitance than assumed and an overestimate of the mobility. Being an amorphous polymer, PTAA is an excellent platform to study effects relating to the gate dielectrics. The performance of PTAA was greatly enhanced by the use of nonpolar gate insulators and eventually equal mobility was achieved in OFETs to that of the bulk mobility determined by transient photoconductivity. It was argued that polar polymers increase carrier trapping and disorder effects in the interface layer. Nonpolar (low-k) organic insulators performed consistently better in both the top and bottom gate devices [26]. Besides PTAA, fluorene–thiophene copolymers, P3HT, polyphenylenevinylenes, and even molecularly doped polymers showed increased mobility with low-polarity resins [131]. It was suggested that the mobility increased because of reduced dipolar (energetic) disorder at the interface, and the best results were obtained when the dielectric layer had a permittivity less than 2.2 and was homogenous (Figure 7.23). The difference in energetic disorder could be gauged by the reduced thermal activation of the field effect mobility with low-k insulators. Lower surface energy dielectrics (organic or inorganic) resulted in progressive increase of the mobility in both PTAA and P3HT (Figure 7.23, polar interface groups lead to broadening of the DOS or create
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Conjugated Polymer Electronics—Engineering Materials and Devices
*
n CH3
CYTOPTM
*
CF2
*
CF
CF
O *
n
*
O
O
O
n*
CF2
O
Polypropylene
CF2
Polyethylene terephthalate (PET)
CF2
CH3 CH3 * n
n
*
*
Polypropylene-co-1-butene
H
n
*
O H H
* n
OR H
Cyanoresin
O Si
CH3
Si
OH Polyvinylphenol
n Si
CH3
X = H, Cl X poly-p-xylylene
n*
R = CH2CH2CN or H
PMMA O
Poly-α-methylstyrene
*
*
* n
O
RO
CH3 *
H
O
* n
RO
*
Polyvinylalcohol
H3C *
*
n OH
Si O
* CH3
n
* BCB
Polyisobutylene
FIGURE 7.22
Organic dielectrics.
localized interface states, which act as traps [26] Figure 7.24). Importantly, the surface energy is only an indicative parameter if the surface is chemically uniform. Organic insulators provide a better-defined interface compared to the complex surface state chemistry of inorganics. It was argued that a thin low
Dipolar disorder at interface
E Bulk N(E)
Drain
Source
Organic semiconductor
Dielectric Gate
FIGURE 7.23 Polar interface groups lead to broadening of the DOS or create localized interface states, which act as traps. (From Veres, J., Ogier, S.D., Leeming, S.W., Cupertino, D.C., and Khaffaf, S.M., Adv. Funct. Mater., 13, 199, 2003. With permission. Copyright 2003, John Wiley & Sons, Inc.)
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10–2
Mobility [cm2 V–1s–1]
Mobility [cm2 V–1 s–1]
10–1
10–2
10–3
10–4
10–3
10–4 Organic Dielectric Inorganic Dielectric
10–5 (a)
0
20
40
60 80 100 Contact angle
120
10–5 (b)
0
20
40 60 80 Contact angle
100
FIGURE 7.24 Field effect mobility in bottom gate P3HT (a) and PTAA (b) devices as a function of water contact angle on the gate insulator. The insulator was either SiO2 with a variety of surface treatments or commercial organic photoresists. (From Veres, J., Ogier, S., Lloyd, G., and de Leeuw, D., Chem. Mater., 16, 4543, 2004. With permission. Copyright 2004, American Chemical Society.)
polarity dielectric is more effective and reliable in providing an optimal interface than a SAM treated SiO2 or Al2O3. Being thicker, (several 10 nm at least) a low-k polymer coating provides better screening from polar defects. Low-polarity interfaces are also useful in reducing the equilibrium water content at the active interface. Large area, low cost TFTs will be required to operate in ambient environment, without expensive encapsulation. Instead of using a high permittivity insulator, the gate capacitance can be increased by an ultrathin SAM layer as dielectric, such as OTS [132]. Self-assembled multilayers have also been demonstrated via selflimiting sequential deposition of siloxane blocks and polarizable stilbazolium layers [133]. Although these approaches are very interesting, the main challenge will be to provide defect-free, robust layers on large areas. Practical devices would greatly benefit from cross-linked insulators for improved layer integrity. Examples of such layers are cross-linked poly(vinyl phenol) (PVP) [134], silsesquioxanes [135], and blends of siloxane precursors with polystyrene or PVP [126]. Recently, Chua et al. showed that benzocyclobutene (BCB, Figure 7.22) was effective for top gate devices with TFB as the channel material [125]. The authors mixed TFB and the liquid monomer BCB and spin-coated the blend solution. In the resulting layer, the BCB monomer is separated to the top and could be cross-linked into a uniform gate dielectric. The ability to cross-link the polymer on top of the OSC without any damage is a remarkable achievement. BCB is also a low-k resin (« ¼ 2.65) and as such provides ideal interface for an organic semiconductor. More recently, the use of BCB enabled the demonstration of ambipolar characteristics of a wide range of organic semiconductors [136].
7.3.10 Ambipolar Materials and Devices Most conjugated materials exhibit p-type behavior, which means that externally injected holes are mobile, whereas electron currents are difficult to observe. Some devices such as OLEDs require both electron and hole conductivity so that the carriers meet and recombine, away from the electrodes. The p- or n-type character of polymers can be probed by steady state current–voltage measurements using different electrodes for carrier injection. Even in otherwise efficient electroluminescent polymers the electron current is orders of magnitude lower than the hole current. It is also much more difficult to gauge the mobility for electrons by transient photoconductivity than for holes. Electron photocurrents are dispersive, indicating strong trapping or very low electron mobility. The p- or n-type character of materials is affected by their chemistry and the degree of overlap between the HOMO and lowest occupied molecular ortbital (LUMO) levels. The molecular relaxation for holes and electrons is also a defining factor. However, it has been difficult to measure intrinsic
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7-25
electron transport properties because impurities and oxygen act as a trap for LUMO states. This means that the carrier range (the product of mobility and lifetime) is too short for electrons to traverse the device without trapping. There are some materials with LUMO levels sufficiently low lying for efficient electron transport, without oxygen becoming a trap. Examples are substituted naphthalene tetracarboxylic diimide (NTDI) [137] and bathophenanthroline [138] derivatives and importantly a ladder-type polymer, poly(benzobisimidazobenzophenanthroline) (BBL) [139]. Classic n-type materials such as C60 and Alq3 exhibit marked differences when studied in air or nitrogen atmosphere. Blends of efficient pand n-type materials have been shown to work in ambipolar devices, i.e., they transport both signs of carriers [140]. Interestingly, a recent work on OFETs demonstrated that it is possible to induce electron currents in a range of polymers that previously were thought of only as p-type semiconductors [136]. Among these were F8T2, P3HT, and PPVs shown in Figure 7.14. The key to this was to use a nonpolar gate dielectric BCB, which was free of hydroxyl groups that also act as electron traps. In addition, the devices were prepared under nitrogen atmosphere and carefully encapsulated. The field effect experiment might have also helped to reveal n-type character because in this configuration a high carrier density is induced in the channel region filling traps. These traps might otherwise mask transport altogether in a small signal regime. It has also been shown that the injected holes and electrons recombine in the channel and the position of the recombination zone can be tuned by the gate voltage [141]. To inject both holes and electrons typically, different metals are used for the source and drain electrodes. The possibility of building ambipolar devices with organic polymers opens the route to light-emitting transistors and CMOS logic. However, there is a serious challenge to develop materials that carry both signs of carriers without the sensitivity to either typical electron or hole trapping impurities. In the meantime, our understanding of n- and p-type conduction mechanisms will be greatly enhanced by investigations of the extent of ambipolar behavior in different materials.
7.4
Expanding the Frontiers of Organic Electronics Using Sequence Independent Synthesis Tools
As was described in the previous sections, much of the progress in the field of organic electronics is still material driven. The inherent versatility of organic chemistry and the potential application of p-conjugated organic materials in structuring molecular electronic components focuses enormous efforts toward the preparation and characterization of new organic-based electronic components [14–17,142–149]. As the field is largely material driven, a major obstacle in organic-based optoelectronics is the difficulty to reach the wide set of material properties required by current device technologies and in synthesizing complicated structures with molecular precision. Using conventional synthetic routes, one is limited by the fact that almost each synthetic step has different and specific conditions and thus, must be studied and optimized. The optimization process is especially important when dealing with multistep syntheses that are a must for preparing materials that are composed of a well-defined sequence of monomers. Furthermore, purification of the intermediates after each synthetic step presents yet another challenge as the molecules become larger. As a result, only a small fraction of the practically limitless choice of organic materials is currently available to the organic semiconductor community. The limited choice of materials consists mainly of noncomplex sequences with purity and molecular precision that are considerably lower than what can be found in biology, where the existence relies mainly on high molecular precision in synthesizing complex molecular entities. In the following text, we will describe synthetic approaches that draw on the methods developed by nature and used in biology.
7.4.1 Nature’s Approach Nature holds two different approaches to the synthesis of biomaterials. Most chemical entities are prepared by target-specific machinery that is tailored to produce one or a very narrow set of structurally similar compounds. This approach is being utilized in most cases where high fidelity in synthesis is
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H H
R1
N PG
H O–
O
+
R2 O
+ H3N O
H PG
H
N PG
R1 H N
O
O H
R2
O
PG + H O 2
SCHEME 7.1 Building a dipeptide from amino acid monomers. PG, protecting group; Rn, the substituent that differ the amino acids from one another.
required, as is the case in the biosynthesis of most small molecules. On the other hand, when versatility is in question, nature adopts a sequence independent synthesis, producing large number of materials that differ in structure and properties out of a very limited number of small building blocks [150]. This approach is being used for the synthesis of polynucleic acids, where the sequence of the monomers in the polymer represents the genetic code, and in peptides, where the sequence of the monomers in the polypeptide chains defines the functionality. Due to the importance of these two reactions to molecular biology, two families of most powerful in-vitro synthetic protocols were developed for the in vitro preparation of peptide and nucleic acid sequences, both in solutions and on solid supports [151,152]. The concept behind these protocols is the use of structurally different building blocks that are capable of interconnecting using the same chemistry and synthetic protocols. For example, peptides are composed of building blocks that interconnect via amide bond formation (Scheme 7.1.) Consequently, any molecule that bears both an amine and a carboxylic acid can be incorporated into a polyamide (polypeptide) in any sequence using the same chemistry for each and every step. An example for such a peptide synthesis is depicted in Scheme 7.2. The coexistence of the two groups that participate in the amide bond formation on the same molecule implies that protection of one of the groups during the coupling process and subsequent deprotection– activation process are also necessary [153] Nevertheless, as the protected groups are always identical, the chemistry of protection–deprotection may also be sequence independent. The recognition that such sequence independent synthetic routes could revolutionize material search, optimization, and even large-scale production led several groups to develop a similar sequence independent approach to the synthesis of p-conjugated oligomers.
7.4.2
Solid-Phase Synthesis of p-Conjugated Molecules
Many solid-phase syntheses that involve the extension of p-conjugated systems are known in the literature. Heck, Stille, Suzuki, Horner–Emmons, Wittig, and other metal-induced coupling reactions were all demonstrated to work on solid supports, some at relatively high yields [154–156]. Table 7.2 summarizes some of these reactions with respect to conditions and yields; many other examples may be found in the literature. Nevertheless, in most cases, the product of such reactions cannot be easily activated to further react in a subsequent step. Additionally, the adoption of solid phase synthesis protocols requires the development of specialized linkers that survive all coupling and deprotection steps without any degradation. The first attempt to develop multistep efficient sequence independent synthesis was reported a decade ago by Moore et al. [162,163]. The group reported the solid-phase synthesis of phenylacetylene oligomers (Scheme 7.3). The synthesis of oligomers utilizes a polymer-bound terminal acetylene that is protected with a trimethylsilyl group. After cleavage of the trimethylsilyl-protecting group and formation of the terminalacetylenic group, a trimethylsiliyl protected aryl iodide was coupled to the polymer-supported terminal acetylene using a palladium(0) catalyst. This cycle was repeated for several times, producing oligomers of varying length and properties. The resulting oligomers were cleaved from the solid support by reacting the triazene group with iodomethane and terminating the oligomer with iodine that could be subsequently reacted in solution to form the trimethyl siliyl–terminated oligomer. The monomers could carry a variety of side groups such as bulky alkyl groups, esters, cyano, and ethers. Both electron accepting and electron donating
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O
H N
O
O O
R1
Piperidine/DMF O NH 2
O
R1
Repeat
R2 HO
N H
fmoc
O HBTU/DIPEA
O O
R1
O
R2
H N O
O
N H
1) Piperidine/DMF 2) TFA/H2O O HO
R1
SCHEME 7.2
R2
H N
NH 2 O
An example of a solid-phase peptide synthesis.
groups showed comparable yields. The reported yield of the reaction was about 50% for the hexamers, significantly lower than the two bioinspired processes. With a similar approach, the group of Anderson was able to produce different unsubstituted oligomers with somewhat lower yields (Scheme 7.4) [164]. Using ortho-, meta-, and para-substituted phenylacetylene monomers, Anderson could prepare oligomers with different molecular geometries, covering different conjugation lengths and molecular shapes. These differences were shown to be translated into a variety of optoelectronic properties. A library of 18 isomeric triisopropylsilyl-capped phenylene ethynylene pentamers was prepared using parallel synthesis, adopting the well-established tea bag approach. Anderson’s group demonstrated the ability to screen the oligomers on the solid support as well as in solutions and, tested selected oligomers for their electroluminescence properties. A different approach was adopted by Ba¨uerle et al. that reported the synthesis of a library of oligothiophene on solid supports (Scheme 7.5 [165,166] and Scheme 7.6 [167]).This group used a standard Merrifield resin or a siliyl group to immobilize the first monomer via an ester linkage. The activation of the thiophene at the 5-position was achieved by site-selective halogenation in the presence of a mercury complex and the extension of the oligomeric skeleton was achieved by Suzuki coupling. The group demonstrated the ability to prepare a 256-member library of quarter thiophene spanning a large spectrum of oxidation potentials, simply by varying the side groups of the monomers in a consistent way. Speicher et al. introduced a somewhat different approach that applied two complimentary oligomer chain elongations steps, Witting and Suzuki coupling, avoiding the use of protecting groups [159]. The group reported a 25% overall yield for an oligomer of four aromatic
O
P1
O
O
O
O
P
OHC
O
O
NH
O O
NH Boc
P O O O
R
O S O
H
NTos
N
OMe OMe
OMe
DMF, NaOCH3 room temp, 48 h
+ – Ph3P Br
Br
DBU, CH2Cl2
O
LHMDS, THF, 25⬚C
P2
O
O
N H
O
O
O
R
O
O S O
O
O
Br
NH N
NHBoc
P1
NH
OMe OMe
OMe
NH-Fmoc
P2
25%b
35%b
100%
a
Yield
[159]
[158]
[157]
Ref.
7-28
Wittig
Horner–Emmons
N H
O
FmocNH
Reaction
Some Approaches to the Coupling of p-Conjugated Systems in the Solid Phase
Horner–Emmons
Reaction Type
TABLE 7.2
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Conjugated Polymers: Processing and Applications
b
Yield of the relevant coupling step. Overall yield after cleavage.
a
Suzuki
Stille
Heck
O N H
H
N
H N
O
O
O
24 h, 1008C
NH2
I
OH B OH
N H
H N
O
908C, 2 h,
Pd(PPh3)4, Na2CO3
Cl
O
O
NaOAc,Bu4NBr, [Pd(OAc)2], DMA,
O
AsPh3, Pd2(dba)3, Dioxane
I
Bu3Sn
pyridine in methanol O2, 2 h
Cu(OAc)2
O
N
H
I
O
O
O
N
H
O
H
N
O
O
NH2
n-PrNH2,O2
Cu(OAc)2
O
Cl
O
65%a
86%b
96%b
[161]
[160]
[160]
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SiMe3 N
N
N
R
R Pd2(dba)3, CuI, PPh3, NEt3
TBAF, THF
I SiMe3
H n = 1–6 R = H, tBu
N
N
N
R
SCHEME 7.3
Solid-phase synthesis of phenylacetylene oligomers.
N N N X = SiMe3 X=H
X TBAF, THF
I
SiMe3
N N N
SiMe3
Y n X = SiMe3 X=H
I SiMe3
TBAF, THF
I Pd2(dba)3, CuI, PPh3, NEt3
n =1-3
N N Si
N 3
1. MeI Si(iPr)3 2. H Pd2(dba)3, CuI, PPh3, NEt3
Si
Si 3
SCHEME 7.4 The synthesis of a solid phase library of pentamers using ortho-, meta-, and para-substituted phenylacetylene monomers.
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R2
R3
R1
R1
R1
O OH Cl
Si
O
Imidazole, DMF
S
Y=H Y=I
Si
S
S
O
Si
O
Y
B O
S
Pd(PPh3)4, THF/H2O, NaHCO3
R4
Pd(PPh3)4, THF/H2O, NaHCO3
R3
R1
S
R2 1) Hg(OCOC5H11)2, CH2Cl2 2) I2, CH2Cl2
Y=H Y=I
1) LDA, THF, −60⬚C 2) I2
R3
R1
S
B O
R3
R1
O S O
Si
S
S
B O
Y
S
S
S O
Si
S
S
R2
R4
S
S
10% TFA, CH2Cl2
S
S
R2
R4
Pd(PPh3)4, THF/H2O, NaHCO3 R2 Y=H Y=I
1) Hg(OCOC5H11)2, CH2Cl2 2) I2, CH2Cl2
SCHEME 7.5
Solid-phase synthesis of thiophene derivatives.
rings (Scheme 7.7). Despite the rather low reported yield, this approach seems very attractive because only chain elongating steps are in use. A symmetric oligomer chain elongation approach was introduced by Tour et al. [168–170]. This group developed a bidirectional approach to the solid phase synthesis of oligo phenylene ethynylenes (Scheme 7.8).
7.4.3 Electronic Peptides Despite considerable efforts, the progress toward a generic sequence independent synthesis of p-conjugated oligomers and polymers, it is evident that the biomimetic processes of peptide and nucleic acid syntheses present superior possibilities and flexibility in making new materials as well as parallel synthesis possible.
R S
O C6H13
C6H13
HO
S
Cs2CO3/K
O
Cl
S
O
O
R
S
S
S
1) I2/Hg(c6H11O2)2 2) CsF, Pd(PP3)4
O
C6H13
C6H13
C6H13 O
S S O
H n+1
C6H13 n = 1-5
SCHEME 7.6
B
1) Bu4NOH 2) MeI
C6H13 HO
S S
O C6H13
H n+1 n = 1-5
Solid phase synthesis of regioregular, head-to-tail coupled, oligo 3-hexyl-thiophenes.
n
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O O
B(OH)2
Br
O
O
O
O
O
O O
O
O Ph O
P+ Ph Ph Br
O
O
O O
B(OH)2 Br
O PG
O O
O PG O
O
O
O
O
SCHEME 7.7
Solid phase synthesis of a p-conjugated oligomer using two complimentary chain-elongating steps.
The basic structure of the building blocks of nucleic acids renders them inherently of no use for direct electronic properties. In contrast, the amide bond exhibits an important contribution of a double bond character, especially when it is electronically coupled to aromatic moieties [171,172]. In most natural amino acids, the p-conjugation along the peptide skeleton is disrupted by the presence of a sp3 carbon atom in between the amine and the carboxy groups. Nevertheless, this carbon atom may be replaced with systems that are p-conjugated (Scheme 7.9). In such a case, the partial double bond character of the amide bond makes it an interesting candidate for coupling conjugated monomers into oligomers and polymers without destroying the p character of the p-conjugated skeleton. Using such a single linker unit will enable the construction of large molecules or polymers with arbitrary, yet well-defined, sequence of monomer units and branching points as well as the tuning of different material properties such as solubility and film-forming ability. This motivated us to explore the use of p-conjugated amino acids as building blocks of p-conjugated oligomers and polymers, with a broader aim of developing massively parallel methods for making and screening complex molecular systems. Ultimately, this approach is expected to allow the physicist or device engineer to specify the required material properties by stating the desired sequence of functional monomers. The synthesis of the different oligo-amides involves only two steps. The first step is the deprotection step, which includes either the transformation of the nitro into the amino group or the removal of a more conventional protecting group such as Fmoc or t-Boc. The second step is the chain elongating step, which is the amide bond formation, using one of several coupling agents such as DCC or HBTU. This simple scheme of high yield and relatively simple deprotection-coupling procedures shows the power of
I
C12H25
C12H25
(CH2)5OH
I
C12H25
C12H25
2
C12H25
I
I
R
C12H25
X = SiMe3 X=H
C12H25
C12H25
(CH2)5OH
C12H25
(CH2)5OH
C12H25
TBAF, THF
(CH2)5OH
(CH2)5OH
(CH2)5OH
X = SiMe3 X=H
(CH2)5OH
C12H25
C12H25
R
The solid phase approach to bidirectional synthesis of oligo 1,4-phenylene ethynylenes.
C12H25
(CH2)5OH
PPTS, n-C4H9OH
C12H25
C12H25
PPh3, Et2NH/THF
Pd2(dba)3, CuI,
SiMe3 SiMe3 Pd2(dba)3, CuI, PPh3, Et2NH/THF
PPh3, Et2NH/THF
SCHEME 7.8
I
C12H25
I
Pd2(dba)3, CuI,
I
I
(CH2)5OH
C12H25
C12H25
I
SiMe3
2 C12H25
SiMe3
(CH2)5OH
TBAF, THF
C12H25
(CH2)5OH
C12H25
C12H25
I
Pd2(dba)3, CuI, PPh3, Et2NH/THF
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R
H O−
O−
H3N+
H3N+
S O
O
SCHEME 7.9 Left hand side—the skeleton of a natural amino acid. Right hand side—an artificial p-conjugated amino acid building block.
the concept of using well-known chemistry of amide bond formation to the synthesis of p-conjugated oligomers. A simple example of such solution phase syntheses is depicted in Scheme 7.10. The absorption and emission spectra of the resulting oligomers clearly indicate that amide bonds are capable of coupling p-conjugated systems. Figure 7.25b depicts the absorption spectra of oligo thiophenes amide, the structure of which is shown in Figure 7.25a. The spectra clearly show the progressive red shift, which indicates an extended conjugation. The charge transport properties of a similar structure were studied using a FET in the bottom contact configuration in which the source and drain contacts are gold electrodes. The chemical structure of the nitro-tri(2,5-thiophene) amide is shown in Figure 7.26a and the schematic of the FET structure is shown in Figure 7.26b. The channel length was varied between 2 and 32 mm and the width was fixed at 6000 mm (COX 43 nF cm1). Extra care was given to remove any residual effects, resulting from the device structure that may interfere with the device material analysis. Figure 7.27a shows the drain-source current as a function of the gate bias for a given drain-source bias of 10 V. In this device, the current is enhanced at positive
C8H17
H
O S
N
nN
C6H13 O
H
(a) (1) Monomer Dimer Trimer Tetramer
(b)
1.0
0.6 3.6
1 2
0.4
Eabs(eV)
Absorption
0.8
3.4 3.2
0.2 3.0
3 4 0.2 0.4 0.6 0.8 1.0 Inverse chain length (1/n)
0.0 300 (b)
400
500
600
Wavelength (nm)
FIGURE 7.25 (a) The structure of t of 2,5-thiophene amide oligomers (n ¼ 1–4) used. (b) The normalized absorption spectra of these oligomers. The inset shows the deduced bandgap as a function of inverse chain length.
S
SCHEME 7.10
H2/Pd
H2N
H2N
H2N
O
N
S
S
S
H
N
O
O
N
H
N
H
O
S
S
S
N
H
H
N
O
O
H O
H2/Pd
S
S
N
H2N
N
H
N
O
O
H
S
N
S
O
H
O
N
H
S
S
S
DCC
O2N
O2N
O2N
N
H
O
S
N H
O
S
O2N
O H
S
O
O2N
S
O H O
O H O
O H O
DCC
O2N
DCC
O2N
S
DCC
O2N
Solution phase synthesis of oligo p-conjugated 2,5-thiopnene amide oligomers.
H2/Pd
H2/Pd
O2N
H
S
S
N H
O
N H
O
S
O
N
H
O
H N
S
O
H
N
S
S
O
H
N H
O
H
N
O
N
S
S
N
O
N
H
O
H
S H
N
O S
O
H N
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H N O2N
S
O
S
O Me
N H
O
S
O
(a) ID
IS
A
Peptide
A
SiO2
A
(b)
VG
FIGURE 7.26 (a) Chemical structure of the nitro-tri(2,5-thiophene) amide used in the FET. (b) A schematic presentation of the bottom contact FET structure.
3.5
3.5
3
3
2.5
2.5
V08 = 20
ID (nA)
ID (nA)
gate-source voltage, indicating that the mobile charges are electrons. Figure 7.27 also shows that at zero gate-bias there is still a significant current flowing between the drain and the source. This may originate from one of the two effects: (a) The effective threshold current is at about 10V. (b) There is bulk conductivity due to the slight charge-transfer nature of the compound, originating from the presence of both electron accepting and electron donating groups on the same p-skeleton. The thiophene-based oligomers were not expected to show good luminescence properties hence, to test for the photoluminescence of such electronic peptides a PPV-like oligomer was synthesized. The chemical structure of a PPV-like oligo peptide is shown in Figure 7.28a and the photoluminescence spectrum of this compound is shown in Figure 7.28b. Using integrating sphere measurements [173] we determined the PL efficiency to be 25%. The optical and charge transport properties of other peptide based p-conjugated systems are currently being explored in our laboratories [174].
2 1.5
VDS = 10V
V08 = 15
2
V08 = 10
1.5
1
V08 = 5
1
0.5
V08 = 0
0.5
V08 = –5 V08 = –10
–30 (a)
–20
–10
10 0 V03 (V)
20
30
0 (b)
5
10 15 V08 (V)
20
25
FIGURE 7.27 Characteristics of the FET with tripeptide 2 as the active layer. (a) The drain-source current as a function of the gate bias for a given drain-source bias of 10 V. (b) The drain current as a function of the sourcedrain bias.
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O O H2N
O O
N O
H
O O
Photoluminescence
(a)
350
400
(b)
450
500
550
600
Wavelength (nm)
FIGURE 7.28 (a) Chemical structure of a PPV-like oligo peptide (b) PL spectrum of the oligo peptide. The quantum efficiency was determined to be 25%.
7.5
Summary and Outlook
In this contribution, we have emphasized the importance of applying an engineering approach to device physics and chemistry as an important ingredient supporting technological progress. We have highlighted that in organic devices, a range of rich and important physical phenomena tend to intermix. Device properties are rarely a straightforward translation of material parameters and they are much influenced by the device configuration and measurement techniques. To describe transport phenomena correctly, device-oriented models are required. Models for electronic transport are often difficult to use across different device platforms. The interplay between device configuration and materials is very much apparent in OFETs, where contacts and interfaces play huge role. The way materials are processed and devices are operated has a strong impact on an experiment. This is particularly an issue due to the sensitivity of organic materials to extrinsic factors. Significant progress has been made on how different materials, for example semiconductors and dielectrics, interact in devices. The use of innovative material combinations will be an interesting route to develop further organic electronics devices. Recent work on ambipolar devices in both blends and single polymers will aid the understanding of n- and p-type transport. Much of this field is still material driven and relies on novel polymeric compounds. As new materials are being constantly developed, it is very important to understand device operation and processing issues and address them via material design. Combining functionality and processability effectively will be critical for new generation polymers. There is a possibility to build materials via new techniques that make the synthesis of functional polymers similar to the engineering of device structures. Solid phase synthetic approaches offer high degree of control over assembly by functionality. We have also described a new direction for semiconductor polymers that utilize peptide chemistry. Electronic peptides may be built upon the well-established peptide synthesis and make the synthesis almost an engineering task.
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145. 146. 147. 148. 149.
Berlin, Y.A., A.L. Burin, and M.A. Ratner. 2000. Superlattices Microstructures 28:241–252. Mathur, N. 2002. Nature 419:573. Willner, I., and E. Katz. 2000. Ang Chem Int Ed 39:1180–1218. Carroll, R.L., and C.B. Gorman. 2002. Ang Chem Int Ed 41:4379–4400. GoldhaberGordon, D., M.S. Montemerlo, J.C. Love, G.J. Opiteck, and J.C. Ellenbogen. 1997. Proc IEEE 85:521–540. Stryer, L. 1996. Biochemistry, 4th ed. Heidelberg: Spektrum. Bodanszky, M., and A. Bodanszky. 1994. The practice of peptide synthesis, 2nd ed. Berlin: Springer. Townsend, L.B., and R.S. Tipson. 1986. Nucleic acid chemistry: Improved and new synthetic procedures, methods, and techniques. New York: John Wiley and Sons. Chan, W.C., and P.D. White. 2000. Fmoe Solid Phase Peptidle Synthesis—A Practical Approach. Oxford: Oxford University Press. Booth, S., P.H.H. Hermkens, H.C.J. Ottenheijm, and D.C. Rees. 1998. Tetrahedron 54:15385–15443. Hermkens, P.H.H., H.C.J. Ottenheijm, and D.C. Rees. 1997. Tetrahedron 53:5643–5678. Hermkens, P.H.H., H.C.J. Ottenheijm, and D. Rees. 1996. Tetrahedron 52:4527–4554. Wang, G., and S.Q. Yao. 2003. Org Lett 5:4437–4440. Couladouros, E.A., and A.D. Magos. 2005. Molecular diversity 9:111–21. Speicher, A., T. Backes, and S. Grosse. 2005. Tetrahedron 61:11692–11696. Stieber, F., U. Grether, and H. Waldmann. 2003. Chem Eur J 9:3270–3281. Colombo, A., J.C. Fernandez, N. de la Figuera, D. Fernandez-Forner, P. Forns, and F. Albericio. 2005. QSAR Comb Sci 24:913–922. Young, J.K., J.C. Nelson, and J.S. Moore. 1994. J Am Chem Soc 116:10841–10842. Nelson, J.C., J.K. Young, and J.S. Moore. J Org Chem 61:8160–8168. Anderson, S. 2001. Chem Eur J 7:4706–4714. Briehn, C.A., M.S. Schiedel, E.M. Bonsen, W. Schuhmann, and P. Bauerle. 2001. Ang Chem Int Ed 40;4680. Briehn, C.A., and P. Bauerle. 2002. J Comb Chem 4:457–469. Kirschbaum, T., and P. Bauerle. 2001. Synth Met 119:127–128. Huang, S.L., and J.M. Tour. 1999. Tetrahedron Lett 40, 3347–3350. Huang, S.L., and J.M. Tour. 1999. J Org Chem 64:8898–8906. Huang, S.L., and J.M. Tour. 1999. J Am Chem Soc 121:4908–4909. Kieninger, M., and S. Suhai. 1996. J Mol Struct 375:181–188. Hagler, A.T., L. Leiserowitz, and M. Tuval. 1976. J Am Chem Soc 98:4600–4612. deMello, J.C., H.F. Wittmann, and R.H. Friend. 1997. Adv Mater 9:230. Tal, S. 2006. In preparation.
150. 151. 152. 153. 154. 155. 156. 157. 158. 159. 160. 161. 162. 163. 164. 165. 166. 167. 168. 169. 170. 171. 172. 173. 174.
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8 Electrical Bistable Polymer Films and Their Applications in Memory Devices 8.1 8.2
Jianyong Ouyang, Chih-Wei Chu, Ricky J. Tseng, Ankita Prakash, and Yang Yang
8.1
Introduction......................................................................... 8-1 Memory Device Using a Blended Film of Polymer and Gold Nanoparticles...................................................... 8-2 Device Using Gold Nanoparticles Capped with Saturated Alkanethiol . Write-Once-Read-Many Times Memory Devices
8.3 8.4 8.5
Memory Device Using Organic Donor and Organic Acceptor............................................................... 8-12 Device Using Composite of Polyaniline Nanofiber=Gold Nanoparticles ........................................ 8-14 Conclusion ......................................................................... 8-20
Introduction
Conjugated polymers and organic molecules are uniquely suitable for thin-film, large area, mechanically flexible, and low-cost electronic devices. Their tremendous potential towards commercial applications has incited a flurry of research, particularly on organic light-emitting diodes [1–3], solar cells [4,5], and transistors [6–8]. On the other hand, nanoparticle is a very different class of material, which is often associated with high-density electronic devices with superior performance and manufacturability [9,10]. Therefore, a conjugated polymer combined with metallic nanoparticles provides an exciting system to investigate the possibility of exhibiting novel functionality. The unique electronic properties of conjugated organic molecules and polymers have originated from their conjugated p-electrons. Similar to inorganic semiconductors, the conductivity of conjugated polymers can be tuned by several orders of magnitude. This can be achieved by a dynamic process such as photoinduced electron-transition process, or electric field-driven charge injection through metal contacts. These processes involve instantaneous charge carriers, which disappear on removal of the external stimulating conditions, such as electric field or light. The conductivity can also be tuned permanently through a chemical oxidation or reduction of the conjugated polymer [11,12]. After oxidation or reduction, the polymer chain is positively or negatively charged, and the polymer film can exhibit high intrinsic conductivity. These properties render many applications to conjugated polymers. Recently, a novel method to tune the conductivity of conjugated polymers or conjugated organic compounds has been proposed and demonstrated in our laboratory [13]. The conductivity is tuned by a dynamic process, like an external electric field, while the conductivity is intrinsic and the charge carriers 8-1
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remain in the system even after removal of the external field. This is achieved on a material system by combining conjugated polymer and metallic nanoparticles into one composite system. Upon the application of external bias, charge transfer process occurs: the metal nanoparticles accept electrons from the conjugated polymer or conjugated organic compound. Therefore, the conjugated polymer or organic compound is positively charged and exhibits a high conductivity. The polymer film is stable in the high-conductivity state for a prolonged period of time, even after removal of the external electric field. The stored charges in the metal nanoparticles are stable due to the insulator coating on the particles. The separated charges may recombine under a reverse biased condition, or sometimes, even by a higher field, and the polymer returns to its low-conductivity state. These processes result in electrical bistability of the polymer film. The conductivity of the polymer film can be programmed by an external electric field, so that the device has strong potential towards application as a nonvolatile memory device. The polymer or nanoparticle memory device is distinctively different from the previous memory devices using organic or polymeric materials, such as organic memory device with a triple-layer structure [14–17], organic memory device with a single organic layer [18–21], write-once-read-many times memory using conducting polymer [22], and memory device using ferroelectric polymer [23]. It has the advantages of a simple device structure, simple fabrication process through a solution processing, and capability to write or erase many times. In this chapter, we review the recent development in polymer memory devices [13,24–27] which have a metal–polymer þ metal nanoparticles–metal structure. Three kinds of material systems for the bistable device were invented. The first one uses a blended film of polymer and metal nanoparticles as the active layer and is presented in Section 8.2. The mechanism for the electrical bistability and memory effect is attributed to an electric field-induced charge transfer between the conjugated structure and the metal nanoparticles. The bistable behavior of the device using organic donor and organic acceptor as the active materials provides direct evidence for this mechanism. This device using organic donor and acceptor is presented in Section 8.3. Finally, a bistable device using a composite with metal nanoparticles directly bonded to the polyaniline (PANI) nanofibers as the active material is discussed in Section 8.4.
8.2
Memory Device Using a Blended Film of Polymer and Gold Nanoparticles [13,24,25]
The chemical structure of some materials used in the memory devices are listed in Figure 8.1. The device has a simple architecture with a blended film of polymer and gold nanoparticles sandwiched between two aluminum (Al) electrodes (Figure 8.2). The electrical behavior of the device strongly depends on the capping thiol molecules on the gold nanoparticles. When gold nanoparticles capped with saturated alkanethiols is used, the device exhibits two electrical states and can be programmed between these two states. Hence, it can be used as a nonvolatile memory device. On the other hand, when gold nanoparticles capped with aromatic thiols, the device cannot be returned to the low-conductivity state after it transits to the high-conductivity state. Thus, this device can be used as a write-once-read-many times memory device.
8.2.1
Device Using Gold Nanoparticles Capped with Saturated Alkanethiol
The active layer of the device is a polymer film consisting of small conjugated organic compound and metal nanoparticle capped with saturated alkanethiol. The gold nanoparticles capped with 1-dodecanethiol (Au-DT NP) (prepared by the two-phase arrested growth method [28]) had a narrow size distribution (1.6–4.4 nm in diameter) and an average particle size of 2.8 nm (Figure 8.3). The alkanethiol is attached to the gold nanoparticle through an Au–S bond. This alkanethiol coating on the gold nanoparticle prevents the aggregation of the gold nanoparticles and gives good solubility in various organic solvents. The device was fabricated through the following process. First, the bottom Al electrode was thermally evaporated on a glass substrate in a very clean chamber under a vacuum of 105 Torr. Then, the active layer between the two Al electrodes was formed by spin coating a 1,2-dichlorobenzene solution of 0.4%
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n
8-3
n O O (b)
(a)
SH N OH
(d)
(c) H | N
H | N
N
N
n
SH (f)
(e) S
S
S
S
(g)
OCH3 O
(h)
FIGURE 8.1 Chemical structure of (a) polystyrene (PS), (b) poly(methyl methacrylate) (PMMA), (c) 8-hydroxyquinoline (8HQ), (d) 1-dodecanethiol (DT), (e) 2-naphthalenethiol (2NT), (f) polyaniline (PANI), (g) tetrathiafulvalene (TTF), and (h) methanofullerene [6,6]-phenyl C61-butyric acid methyl ester (PCBM).
by weight Au-DT NP, 0.4% by weight 8-hydroxyquinoline (8HQ), and 1.2% by weight polystyrene (PS). This polymer film had a thickness of about 50 nm. Finally, the device was completed by thermal evaporation of the top Al electrode. The top and bottom Al electrodes had a line width of 0.2 mm and were aligned perpendicular to each other, so that the device had an area of 0.2 0.2 mm2. This device is represented as Al=Au-DT NPþ 8HQ þ PS=Al in this chapter. Figure 8.4 shows the current–voltage (I– V) curves of Al=Au-DT NP þ 8HQ þ PS=Al, tested in vacuum using an HP V Au nanoparticles Al 4155B semiconductor parameter analyzer. The pristine device exhibited very low current, approximately 1011 A at 1 V. An electrical transition took place at 2.8 V with an abrupt current increase from 1011 to 106 A (curve (a)). The device exhibited good stability in this high-conductivity (ON) Al state during the subsequent voltage scan (curve (b)). The high-conductivity state was able to return to the low-conductivity (OFF) state by applying a negative bias as Middle composite layer indicated in curve (c) where the current FIGURE 8.2 Device structure of polymer memory device. suddenly dropped to 1010 A at 1.7 V.
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40
Count
30 20 10 0 1.5
2.0
2.5 3.0 3.5 Diameter (nm)
4.0
4.5
1 494612 2002 AKU X680K 10nm
FIGURE 8.3 Transmission electron microscope (TEM) image and size histogram of Au-DT NP. (From Ouyang, J., Chu, C.-W., Szmanda, C., Ma, L., and Yang, Y., Nat. Mater., 3, 918, 2004. With permission.)
When the device was tested in nitrogen atmosphere or in air, it exhibited similar electrical behavior. The transition voltages and the current for the device in the high-conductivity state are almost the same as those tested in vacuum whereas the current for the device in the low-conductivity state tested in air is higher by one to two orders of magnitude than that tested in vacuum. These results indicate that oxygen and moisture do not play a role for the electrical transitions. Other materials were also used to fabricate the device to study the effect of materials on the device performance. When 8HQ was replaced by other conjugated organic compounds, such as 9,10-dimethylanthracene, PS replaced by poly(methyl methacrylate) (PMMA), or gold nanoparticle replaced by silver nanoparticle, similar electrical behavior was observed. However, when a PMMA film consisted
10−5
Current (A)
(b)
10−7
(a) (c)
10−9
10−11
−2
−1
0
1
2
3
4
5
Bias (V)
FIGURE 8.4 I–V curve of a device Al=Au-DT NP þ 8HQ þ PS=Al. (a), (b), and (c) represent the first, second, and third bias scans, respectively. The arrows indicate the voltage-scanning directions. (From Ouyang, J., Chu, C.-W., Szmanda, C., Ma, L., and Yang, Y., Nat. Mater., 3, 918, 2004. With permission.)
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6
8-5
W
Bias (V)
4 2
R
R
0 E
−2 −4 Current (A)
10−6
1
10−7 10−8 0
10−9 0
1
2
3
4
5
Time (s)
FIGURE 8.5 Write–read–erase cycles of the device Al=Au-DTNP þ 8HQ þ PS=Al. The top and bottom curves are the applied voltage and the corresponding current response, respectively. W, R, and E in the top figure mean write, read, and erase, respectively. ‘‘1’’ and ‘‘0’’ in the bottom figure indicate the device in the high- and low-conductivity state, respectively. (From Ouyang, J., Chu, C.-W., Szmanda, C., Ma, L., and Yang, Y., Nat. Mater., 3, 918, 2004. With permission.)
of only Au-DT NP without 8HQ or any other conjugated organic compound, no remarkable electrical switching was observed. Switching between the high- and low-conductivity states of Al=Au-DT NP þ 8HQ þ PS=Al was performed numerous times. The device was written, read, and erased repeatedly in air, as demonstrated in Figure 8.5 (for convenience the absolute value of the current is shown). A voltage of 5 V was applied to write ‘‘1’’ to the device, that is, this voltage switched the device to the high-conductivity state. (Here we designate ‘‘1’’ as the high-conductivity state and ‘‘0’’ as the low-conductivity state.) This ‘‘1’’ state could be read by a low voltage (1.1 V in our case). The current during the read pulse was in the range of 107 A. This high-conductivity state was erased by a voltage of 2.3 V, which returned the device to the lowconductivity ‘‘0’’ state. This ‘‘0’’ state could also be detected by applying a small voltage. The current at 1.1 V was in the range of 109 A. These write–read–erase cycles demonstrate that the device can be used as a nonvolatile digital memory device. The device in the low-conductivity state can be switched to high conductivity by a pulse of 5 V with a width of 25 ns. The device, which was in the low-conductivity state, exhibited current less than 109 A in the voltage range of 0–1 V. It exhibited a current with four orders of magnitude higher after applying a pulse of 5 V with a width of 25 ns. At this time, our equipment (HP 214B Pulse Generator) was incapable of generating a pulse shorter than 25 ns. It is possible that the response time of this device is actually faster than what we are able to measure. The polymer memory device has a simple device structure and can have very high density when highdensity electrodes are fabricated. The device structure can be even simpler, and the density can be pushed to very high values, when the operation of the device is combined with an atomic force microscope (AFM). A schematic operation configuration for this device is shown in Figure 8.6. The device was fabricated by one step: the polymer film was spin coated on a conductive substrate. The conductive substrate was used as the bottom electrode whereas an AFM tip as the top electrode. Figure 8.7 shows a surface potential AFM picture of an Au-DT NP þ 8HQ þ PS film on Al, deposited on silicon wafer. At first, an area of 20 mm 10 mm of the film was scanned vertically in contact mode while applying þ10 V bias through a 50 nm silicon nitride AFM tip coated with Au. Then, another area of 20 mm 5 mm was scanned horizontally while applying a 10 V bias through the tip. Finally, the scanning surface potential image was taken with a tapping model. A dc bias of 4 V was applied on the
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film through the 50 nm AFM silicon nitride tip coated with Au. The two pretreated areas exhibited remarkably different potential in the surface potential AFM. These AFM tip V experiments demonstrated the feasibility of the device operation combined with an AFM. Polymer film The nanosecond scale transition of these devices and the electric field-induced change of the surface potenBottom electrode tial of the polymer film suggest that the switching process may be an electronic process rather than chemical FIGURE 8.6 Test configuration for the operrearrangement, conformational change, or isomerizaation of the device using an AFM tip as the top tion. For example, conformational change for a ferroelectrode. (From Ouyang, J., Chu, C.W., Tseng, R.J.H., Prakash, A., and Yang, Y., Proc. IEEE, 93, electric polymer occurs on a 30 ms timescale [23]. Atomic movement or molecular isomerization that can 1287, 2005. With permission.) result in electrical bistability is observed on molecular devices [29,30]; however, the transition time is on the millisecond timescale or even longer. We study the switching mechanism by ac impedance spectroscopy, study of the transport mechanisms, and the energy levels of the materials. Our experimental results suggest that the switching mechanism is not due to the formation of conductive filaments between the two metal electrodes, which was observed in a polymer film by others [31,32]. It is unlikely that filament formation is the reason for the electronic transitions in our device, since the electrical behavior of our device is strongly dependent on the structure and concentration of the gold nanoparticles. In addition, ac impedance studies, from 20 to 106 Hz (Figure 8.8), indicate that the electronic transitions in our device are different from the dielectric breakdown found in polymer films. We observed dielectric breakdown in a device with a polystyrene film sandwiched between two Al
10
20
30
Dimension (µm)
FIGURE 8.7 Scanning surface potential AFM image of Au-DT NP þ 8HQ þ PS film with Al as bottom electrode and silicon wafer as substrate. The vertical bar with yellow color was pretreated with a þ 10 V dc bias whereas the horizontal bar with brown color was pretreated with a 10 V dc bias. (From Ouyang, J., Chu, C.-W., Szmanda, C., Ma, L., and Yang, Y., Nat. Mater., 3, 918, 2004. With permission.)
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8-7
Capacitance (F)
Capacitance (F)
electrodes. After breakdown, the current 1.5x10–10 increased by more than four orders of mag(a) nitude, and the capacitance was lowered by about one order of magnitude in the whole 1.0x10–10 frequency range. In comparison, Al=AuDT NP þ 8HQ þ PS=Al exhibited an increase in current by more than four orders of magnitude. The capacitance increased in 5.0x10–11 the low-frequency range whereas remained (b) the same in the high-frequency range after 0.0 a pristine device was turned to the highconductivity state. This suggests that space charges may generate in the film after the as(c) prepared device is electrically turned to a high-conductivity state. Polystyrene may 1.5x10–10 act as an inert matrix for Au-DT nanoparticles and 8HQ, and not play a role in the (d) electronic transition. So, we exclude from our model changes in electrical behavior based on slow speed switching mechanisms and filament formation. 1.0x10–10 103 104 106 105 The conduction mechanism for Al=AuDT NP þ 8HQ þ PS=Al in the low-conFrequency (Hz) ductivity state may be due to quite small FIGURE 8.8 Capacitance of (a) pristine Al=PS=Al, (b) amount of impurity or hot electron injecAl=PS=Al after breakdown, (c) Al=Au-DT NP þ 8HQ tion. The conduction mechanism for the þ PS=Al in high-conductivity state, and (d) pristine Al=Audevice in the high-conductivity state was DT NP þ 8HQ þ PS=Al. (From Ouyang, J., Chu, C.W., Tseng, studied by investigating the temperature R.J.H., Prakash, A., and Yang, Y., Proc. IEEE, 93, 1287, 2005. With permission.) dependence of the current and the analysis of the current–voltage relation. The current for the device in the high-conductivity state was almost temperature-independent (Figure 8.9), and the I–V curves could be fitted well by a combination of direct tunneling (tunneling through a square barrier) and Fowler–Nordheim tunneling (tunneling through a triangular barrier) (Figure 8.10) as given by the following expression [33]: I ¼ C1 V e
ffiffiffiffiffiffiffi
p 2d 2m F h
þ C2 V 2 e
ffiffiffiffiffi
p 4dF3=2 2m 3qhV
The first term on the right-hand side of the equation is the current contributed by direct tunneling, and the second term is the current contributed by Fowler–Nordheim tunneling. In this equation, d is the tunneling distance, m* is the effective mass of the charge carrier, and F is the energy barrier height. At low voltage, V < F, direct tunneling is the dominant conduction mechanism, and at high voltage, V > F, Fowler–Nordheim tunneling becomes the dominant conduction mechanism. The different conduction mechanisms in the two states suggest change in the electronic structure of the device after the electrical transition. It has already been demonstrated that 8HQ and gold nanoparticle can act as electron donor and acceptor, respectively [34–37]. Moreover, the different surface potentials of the Au-DT NP þ 8HQ þ PS film after treatment by different electric fields as shown in Figure 8.7 suggest that electric field can induce polarization of the film. Hence, we propose a charge transfer between Au-DT NP and 8HQ under a high electric field for the electronic transition in Al=AuDT NP þ 8HQ þ PS=Al. Before the electronic transition, there is no interaction between the Au-DT nanoparticle and 8HQ. Concentration of charge carriers due to impurity in the film is quite low, so that
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the film has very low conductivity. However, when the electric field increases to a 4V certain value, electron on the highest occu–12.5 pied molecular orbital (HOMO) of 8HQ 3V may gain enough energy to tunnel through the capped molecule, 1-dodecanethiol, –13.0 2V into the gold nanoparticle (Figure 8.11). Consequently, the HOMO of 8HQ becomes partially filled, and 8HQ and gold –13.5 nanoparticle are charged positively and negatively, respectively. Therefore, carriers 1V are generated and the device exhibits a –14.0 0.006 0.008 0.010 0.004 high-conductivity state after the charge transfer. It is well known that the conduct1/T (1/K) ivity of conjugated organic compounds FIGURE 8.9 Arrhenius plot of the temperature dependence of will increase after their HOMO or lowest current for Al=Au-DT NP þ 8HQ þ PS=Al in the high-conunoccupied molecular orbital (LUMO) ductivity state at applied voltages of 1, 2, 3, and 4 V. (From becomes partially filled [17,38,39]. Ouyang, J., Chu, C.-W., Szmanda, C., Ma, L., and Yang, Y., Nat. For the device in the high-conductivity Mater., 3, 918, 2004. With permission.) state, the charge transport through the polymer film may take place through charge tunneling among the 8HQ molecules. The separation among the 8HQ molecules in the polymer film will become larger than that in the 8HQ crystal. A simple estimation suggests that the separation among the 8HQ molecules in the Au-DT NP þ 8HQ þ PS film is about 6–10 A˚. For such a separation, it is reasonable that the tunneling process becomes the dominant charge transport mechanism among the 8HQ molecules. The simple model can interpret the stability of the device in the high-conductivity state as well as the erasing process by applying a negative bias. Stability of the negative charge on gold nanoparticle is due to the insulator coating, 1-dodecanethiol, on the gold nanoparticles, which prevents recombination of the charge after removal of the external electric field. Since the charge transfer is induced by an external electrical field, the film is 4.0x10–6 polarized after the charge transfer. Only a reverse electric field can assist the tunneling 3.0x10–6 of the electron from the gold nanoparticle back to the HOMO of 8HQþ, resulting in a return to the low-conductivity state. 2.0x10–6 The electric field-induced charge transfer model explains quite well the elec1.0x10–6 tronic transition observed in the device, Al=Au-DT NP þ 8HQ þ PS=Al, and is supported by the evidences noted above 0.0 as well as the following additional evidences: fast switching speed, lack of dielec0 2 4 6 8 10 tric breakdown, temperature insensitive Bias (V) conductivity, and device performance deFIGURE 8.10 I–V curve of the device, Al=Au-DT NP þ 8HQ pendence on the gold nanoparticle concenþ PS=Al, in the high-conductivity state. The scattered points tration and size. Charge tunneling through are the experimental results, the solid line is the data fit comthe insulator coating on the gold nanoparbining direct tunneling and Fowler–Nordheim tunneling, and ticle is possible, as required by our model. the broken line is the data fit of Fowler–Nordheim tunneling. For example, it has been observed fre(From Ouyang, J., Chu, C.-W., Szmanda, C., Ma, L., and Yang, Y., Nat. Mater., 3, 918, 2004. With permission.) quently, by electrochemical measurements, Current (A)
In (I )
–12.0
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e
N OH E
FIGURE 8.11 Schematic electron transfer from 8HQ to the core of the gold nanoparticle. The yellow circle indicates the core of the gold nanoparticle, and the gray ring indicates the capped 1-dodecanethiol. The curved arrow denotes the electron transfer from 8HQ to the core of gold nanoparticle, e denotes the electron. The direction of the electric field (E) is represented by the linear arrow. (From Ouyang, J., Chu, C.W., Tseng, R.J.H., Prakash, A., and Yang, Y., Proc. IEEE, 93, 1287, 2005. With permission.)
d¼
8-9
that gold nanoparticles coated with an insulating alkanethiol layer can be reduced or oxidized [40,41]. This mechanism is further supported by the energy levels of the materials. Figure 8.12 shows the energy levels of gold nanoparticle and 8HQ and the electron transfer from the HOMO level of 8HQ to the core of the gold nanoparticle. The HOMO and LUMO levels of 8HQ, calculated using density function theory using Becke’s three parameter functional with the Lee, Yang, and Lee correlation functional (DFT B3LYP) method with the 6–31 þ G(d,p) basis set [42] are 1.9 and 6.1 eV, respectively. The quantized energy d for a gold nanoparticle with a diameter of 2.8 nm is about 0.014 eV calculated in terms of the equation [43]:
4EF 3N
where EF is the Fermi energy of bulk gold, and N denotes the number of gold atoms in one gold nanoparticle. This quantized energy is much smaller than the thermal energy at room temperature, so that its effect can be neglected. On the other hand, the Coulomb energy (Ec) to charge a gold nanoparticle with a diamE eter of 2.8 nm and capped with 1-dodecanethiol is about 0.1 eV calculated by the LUMO following equation [44]: LUMO
Ec ¼
e2 2C
and Work function HOMO
HOMO Au NP
DT
8HQ
FIGURE 8.12 Energy diagram of the core of gold nanoparticle, 1-dodecanethiol (DT), and 8-hydroxyquinoline (8HQ). The two dots on the HOMO of 8HQ represent two electrons. The linear arrow indicates the direction of the electric field (E), and the curved arrow indicates the electron transfer from 8HQ to the core of gold nanoparticle. (From Ouyang, J., Chu, C.W., Tseng, R.J.H., Prakash, A., and Yang, Y., Proc. IEEE, 93, 1287, 2005. With permission.)
r C ¼ 4p«0 «r (r þ d) d where C is the capacitance of the gold nanoparticle, «0 the permittivity of free space, «r the permittivity of the capped molecule on the gold nanoparticle, r the radius of the gold nanoparticle core, and d the length of the capped molecule. This charging energy is the energy to be overcome for the charge transfer to take place. It is possible for the electron to gain such energy under a high electric field. These
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considerations on the energies further suggest that a charge transfer from 8HQ to gold nanoparticle is possible under the application of a high electric field.
8.2.2
Write-Once-Read-Many Times Memory Devices
Current (A)
The electrical behavior of the polymer memory device was dependent on the chemical structure of the metal nanoparticles in the polymer film. When the chemical structure of the capping molecule on the gold nanoparticle changed from a saturated alkanethiol to an aromatic thiol, the electrical behavior of the device significantly changed. A device using a PS film consisting of gold nanoparticles capped with 2-naphthalenethiol (Au-2NT NP) was fabricated through a similar fabrication process as that of Al=AuDT NP þ 8HQ þ PS=Al. This device is represented by Al=Au-2NT NP þ PS=Al, and its I–V curve is shown in Figure 8.13. In the first voltage scan, the current exhibited rapid increase starting from about 4 V, but this current increase was less abrupt than that in the scan of the device Al=Au-DT NP þ 8HQ þ PS=Al. After the voltage scan from 0 to 8 V, the device transited to a high-conductivity state. The current at 2 V for the device in the two conductivity states was different by about three orders in magnitude. This device after transiting to the high-conductivity state could not return to the lowconductivity state by applying a negative voltage or a high voltage in either polarity. The presence of the other conjugated organic compound, such as 8HQ, in the polymer film did not affect the I–V curve of the device. This is because conjugated structure, which attaches to the gold nanoparticles is already present in the polymer film, and the aromatic group on the capped molecule screens the interaction between the core of the gold nanoparticle and the conjugated compound. The devices using gold nanoparticles capped with aromatic thiols can transit to a high-conductivity state and have good stability in that state. Therefore, these devices can be used as write-once-read-many times memory devices. To further understand the electric field-induced transition, I–V curves in both states were analyzed in terms of different theoretical models [45]. A linear relationship between log I and V1=2 in the voltage range of 0–3 V (Figure 8.14a) suggested that the current before the electrical transition was controlled by charge injection from the Al electrode into the Au-2NT NPs. On the other hand, the I–V relation changed after the electrical transition: log I had a linear relation to log V (Figure 8.14b), and a best fit of log I–log V indicated that I / V1.9. Hence, the current through the device becomes a space-charge-limited current after the transition. The current in the high-conductivity 10–3 state increases with increase in the concentration of the Au nanoparticles in the film 10–4 (Figure 8.15). The current at 6 V was plotted versus the Au-2NT NP concentration in 10–5 Figure 8.16. A best fitting of this data indi10–6 cates that the current is proportional to the square of the concentration. This is 10–7 because both the charge carrier population and the charge carrier mobility are depen–8 10 dent on the concentration of the Au-2NT NP in the polymer film. These results sug10–9 gest that the Au-2NT NP is the media for the charge transport through the film due 0 2 4 6 8 to the conjugated naphthalene structure. Bias (V) This is different from the film of Au-DT FIGURE 8.13 I–V curves of a device Al=Au-2NT NP NP þ 8HQ þ PS, where there is only 8HQ þ PS=Al. The arrows indicate the voltage-scanning directions. with conjugated quinoline structure and (From Ouyang, J., Chu, C.W., Sievers, D., and Yang, Y., Appl. 8HQ is the media for the charge transfer Phys. Lett., 86, 123507, 2005. With permission from American through the film. Institute of Physics.)
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Current (A)
10–9
10–10
10–11
10–12 0.0 (a)
0.5 1.0 Voltage1/2 (V 1/2)
1.5
10–2
Current (A)
10–3
10–4
10–5
10–6 0.1 (b)
1 Voltage (V)
10
FIGURE 8.14 I–V curves of Al=Au-2NT NP þ PS=Al (a) before and (b) after the transition. The scattering of the data at low voltage is due to experimental error at very low current. (From Ouyang, J., Chu, C.W., Sievers, D., and Yang, Y., Appl. Phys. Lett., 86, 123507, 2005. With permission.)
The effect of film thickness on the transition voltage was investigated as well (Figure 8.17). The transition voltage was defined as the voltage at which the current starts a rapid increase in the log I versus V graph. The transition voltage for some devices exhibited scattered values, and all these voltages were plotted in Figure 8.17. The transition voltage increases linearly with the increase of the film thickness. This linear relation indicates that the electrical transition is a result of an electric field effect. The electric field-induced charge transfer in the polymer film is further evidenced by the asymmetric I–V curve for the device in the high-conductivity state. The I–V curve was symmetric in the two polarity directions in the voltage range below 3 V before the transition whereas it became asymmetric after the transition to the high-conductivity state. If the polarity direction that resulted in the transition was defined as positive, the current along the negative direction was higher than along the positive direction after the transition (Figure 8.18). The current at 3 V reached almost 10 times as that at 3 V. This asymmetric I–V curve after the transition suggests that the electric field may induce a polarization of the Au-2NT NP in the polymer film. Based on these experimental results, the authors proposed an electric field-induced charge transfer similar to the device Al=Au-DT NP þ 8HQ þ PS=Al. The polarization of the polymer film after the electronic transition is interpreted as the result of an electric field-induced charge transfer between the gold
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10–2 (a)
Current (A)
10–4
(b) (c)
10–6
10–8
10–10
0
2
4
6 Bias (V)
8
10
12
FIGURE 8.15 I–V curve of Al=Au-2NT NP þ PS=Al with different Au-2NT NP concentrations in the high-conductivity state: 1% (a), 0.4% (b), and 0.1% (c). (From Ouyang, J., Chu, C.W., Tseng, R.J.H., Prakash, A., and Yang, Y., Proc. IEEE, 93, 1287, 2005. With permission.)
8.3
nanoparticle and the capping 2NT: the capping 2NT donates an electron to the core of the gold nanoparticle (Figure 8.19). After the charge transfer, the Au-2NT NP polarizes along the applied electric field. Hence, the device will have higher current when the external electric field is applied along this polarization direction than that when the external electric field is applied against this polarization. The model can explain the two conductivity states well. Before the transition, the current is controlled by the charge injection from the electrode into the polymer film due to a big energy barrier between the Al electrode and Au-2NT NPs. After the charge transfer between the Au nanoparticle and 2NT induced by a high electric field, 2NT is positively charged so that the film exhibits a high current.
Memory Device Using Organic Donor and Organic Acceptor [26]
Current at 6 V (A)
Electrical bistability can also be found in all-organic systems possessing both an electron donor and acceptor. In these systems, the electronic switching mechanism is attributed to an electric field-induced charge transfer between conjugated organic compounds, in fact this system presents one of the most direct proofs of this mechanism. This idea is tested by using methanofullerene [6,6]phenyl C61-butyric acid methyl ester (PCBM) as an organic electron acceptor and tetrathiafulvalene (TTF) as an organic 10–4 electron donor. The active layer of this device is formed by spin coating a 1,2-dichlorobenze solution with 1.2% by weight of polystyrene, 0.8% by weight TTF, and –5 10 0.8% by weight PCBM. The device exhibits electrical bistable behavior as shown in Figure 8.20. The 10–6 voltage to turn the device from the lowto high-conductivity state is 2.6 V. At this critical voltage, the current increases 10–7 abruptly from 107 to 104 A, and the 0.1 1 device is then stable in this high-conductAu NP concentration (%) ivity state. The low-conductivity state can FIGURE 8.16 Dependence of current at 6 Von the Au-2NT NP be recovered by simply applying a higher concentration for a device Al=Au-2NT NP þ PS=Al in the highvoltage bias at either polarity (e.g., 9 or conductivity state. The scatters are the experimental data, and 9 V can be used) as indicated by curve the line is a best linear fitting on the data. (From Ouyang, J., Chu, (c), where the current suddenly drops C.W., Tseng, R.J.H., Prakash, A., and Yang, Y., Proc. IEEE, 93, from 104 to 106 A at 6.5 V. The 1287, 2005. With permission.)
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Current (A)
VT (V)
switching time for the device from ON to OFF or from OFF to ON state is 10 shorter than 100 ns. The stability of the devices is first 8 tested through stress testing of the device. For this experiment, a constant voltage of 0.5 V is applied to the devices 6 in both the OFF and ON states and the current output versus time is recorded. 4 It is noted that there is no significant current change for the devices in either state even after 12 h of continuous 2 stress testing. Stability can also be evaluated by determining the retention 0 20 40 60 80 100 120 ability, which is measured by leaving Thickness (nm) several devices in the high-conductivity FIGURE 8.17 Dependence of transition voltage on the film state without any applied voltage bias thickness for a device Al=Au-2NT NP þ PS=Al. The scatters are under nitrogen environment. It is witthe experimental data, and the line is a best linear fitting on the nessed that once the device is in the ON data. (From Ouyang, J., Chu, C.W., Tseng, R.J.H., Prakash, A., and state, it will remain that way for several Yang, Y., Proc. IEEE, 93, 1287, 2005. With permission.) days, sometimes even for weeks. This device can also be successfully cycled between the two states many times, allowing implementation into rewritable memory applications. To understand the electrical switching mechanism, we study the conduction mechanism of the devices in both states. Before the electrical transition, there was a linear relationship between log I and V1=2 in the voltage range from 0 to 1.7 V (Figure 8.21a). Such linearity suggests that the conduction mechanism is probably thermionic emission [46], that is, the conduction mechanism is dominated by charge injection. After the electrical transition, a linear relationship is observed for a log (I=V) versus V1=2 (Figure 8.21b). Hence, Poole–Frenkel emission (P–F) is probably the conduction mechanism for the device in the high-conductivity state. The P–F conduction mechanism is dominated by charge transport through the bulk material, which are filled with electrically charged defects [47,48]. The presence of this P–F mechanism is further confirmed by using electrodes of dissimilar work functions, i.e., with the ITO=PS þ PCBM þ TTF=Al configuration, by symmetric I–V 0.0 characteristics for both polarities. Therefore, the current conduction changes from an injection-dominated mechanism −2.0 3 10−6 in the OFF state to charge transport dominated mechanism in the ON state. The switching mechanism is further −4.0 3 10−6 studied using ac impedance spectroscopy. The device in the OFF state exhibits a −6.0 3 10−6 capacitance of 30 pF from 20 to 106 Hz unaffected by changes in frequency, how0 1 2 3 −3 −2 −1 ever after the device is switched to the Bias (V) ON state the capacitance becomes strongly dependent on the frequency FIGURE 8.18 Asymmetric I–V curve of Al=Au-2NT NP (Figure 8.22). For devices in the ON þ PS=Al in the high-conductivity state. (From Ouyang, J., state, the capacitance in the highChu, C.W., Sievers, D., and Yang, Y., Appl. Phys. Lett., 86, 123507, 2005. With permission.) frequency range of 104106 Hz is almost
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the same as in the OFF state, then it subsequently increases with decreasing frequency. The capacitance in ON state is higher than in OFF state by more than one order of magnitude at the e low frequency less than 600 Hz. This S increased capacitance at low frequency may be influenced by the increase of the apparent dielectric constant of the film. The change of apparent dielectric conS S Au NP stant is believed to be associated with the field-induced dipole formation between the donor and the acceptor. S In light of the above experimental results and discussions, we propose that the electronic transition is due to an electric field-induced charge transfer in the film between TTF and PCBM. This can be partially confirmed by observing that the bistable phenomenon e does not exist if only one component is FIGURE 8.19 Schematic electron transfer from capping molpresent. By noting that the HOMO and ecule 2-naphthalenethiol to the core of the gold nanoparticle. LUMO levels for TTF are 5.09 and 2.33 Only four 2NT molecules are plotted to represent all-capping eV [49] and for PCBM they are 6.1 and 2NT molecules on the gold nanoparticle. The curved arrow indi3.7 eV [50], it can be inferred that TTF cates the electron (e) transfer. The linear arrow denotes the direcand PCBM are electron donor [51,52] tion of the electric field (E). (From Ouyang, J., Chu, C.W., Tseng, and acceptor [53], respectively. The R.J.H., Prakash, A., and Yang, Y., Proc. IEEE, 93, 1287, 2005. With materials are chosen because their enpermission.) ergy levels do not allow electron transfer in their ground states. Also the resulting UV–visible absorption spectrum is a superposition of individual TTF and PCBM spectra. Therefore, the interaction between TTF and PCBM may be weak before the electronic transition. In addition, concentration of charge carriers due to impurity in the film is quite low, so that the film has low conductivity. However, a high electrical field may facilitate an electron transfer from the HOMO of TTF to the LUMO of PCBM. Consequently, the HOMO of TTF and LUMO of PCBM become partially filled, and TTF and PCBM are charged positively and negatively, respectively. Therefore, carriers are generated and the device exhibits sharp increase in conductivity after the charge transfer.
8.4
Device Using Composite of Polyaniline Nanofiber=Gold Nanoparticles [27]
By exploiting the advantages of nanocomposite materials, further progress in the development of polymer electronic memory devices can be made. As discussed above, the electrical bistability is related to the electric field-induced charge transfer between two components. Thus, the compatibility of these two components is crucial for optimal device performance. Enhancing the interface between the two components could lead to superior device performance, so that to this purpose a nanocomposite of polyaniline nanofiber–Au nanoparticle is used as the active material. A conducting polymer, such as polyaniline, decorated with metallic or semiconducting nanoparticles furnishes an exciting system to investigate the prospect of designing device functionality directly into the material. A facile bulk synthesis method capable of producing high-quality polyaniline nanofibers, with
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0.001
0.0001
Current (A)
(c) (b)
10–5
(a) 10–6 10–7 10–8 10–9 –8
–6
–4
–2
0
2
4
Voltage
FIGURE 8.20 I–V curve of a device, Al=PS þ PCBM þ TTF=Al. (a), (b), and (c) represent the first, second, and third bias scans, respectively. The arrows indicate the voltage-scanning directions. (From Chu, C.W., Ouyang, J., Tseng, R.J.H., and Yang, Y., Adv. Mater., 17, 1440, 2005. With permission.)
diameters tunable from 30 to 120 nm, is employed [54–56]. The nanocomposite is prepared by growing gold nanoparticles on the polyaniline nanofibers through the reduction of chloroauric acid (HAuCl4) in an aqueous solution containing the nanofibers [57–59]. A transmission electron microscopic (TEM) image of the PANI nanofiber–gold nanoparticle composite is shown in Figure 8.23. An active film is formed by spin coating an aqueous solution of 0.1% by weight PANI nanofiber–gold nanoparticle composite in 1.5% by weight polyvinyl alcohol, with an overall film thickness of 70 nm. Polyvinyl alcohol serves as an electrically insulating matrix for the nanocomposite. The actual device structure and fabrication process are similar to the methods used to fabricate devices with the Al=Au-DT NP þ 8HQ þ PS=Al system. The PANI nanofiber–gold nanoparticle device can be used as an electronic memory device due to its bistable electrical behavior as shown in Figure 8.24. The voltage used to switch the device from the lowto high-conductivity state is 3 V, where the current transition is from 107 to 104 A. The device is then stable in the high-conductivity state until a reverse bias of 5 V is applied to return the device to the low-conductivity state. When the voltage is raised above 3 V, a region of negative differential resistance (NDR) is observed, but appears to have no effect on the performance of our device within the 3–4 V region. NDR has been reported elsewhere in other memory devices [16,60,61], and here it seems that this NDR behavior is related to the size of the gold nanoparticles. When the gold nanoparticles have diameters greater than 20 nm, the devices can only be switched on once, and during this time in the ON state they exhibit ohmic behavior, indicating that the more metallic nature of larger gold particles dominates the switching. The gold nanoparticles are absolutely critical to the operation of the device, as devices made with only PANI nanofibers do not demonstrate electronic switching. To further demonstrate the feasibility of this material for use as nonvolatile memory, the stability of the system must be evaluated. Data retention is gauged by measuring the current of the device in the ON state over a long period of time; even after a 3-day period no appreciable change in conductivity is observed for these devices although after several days a slight decrease in conductivity in the highconductivity state is observed. This stress test can indicate the device robustness, as it is performed by applying a constant voltage of 1 V and recording the current for every 5 s. In this case, the test was
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log(I/A)
1E –8
InJ −V 1/2 1E –9 Thermonic emission fitting on–off current
(a)
1E –10 0.0
0.4
0.2
0.6
0.8
1.0
1.2
1.4
1.6
V 1/2 5E−5
log(I/V )
4E−5
3E−5
PF fitting 2E−5
0.0 (b)
In(J/V ) −V 1/2
0.5
1.0
1.5
2.0
2.5
V 1/2
FIGURE 8.21 The analysis of I–V characteristic for device Al=PS þ PCBM þ TTF=Al in (a) high- and (b) lowconductivity state. (From Chu, C.W., Ouyang, J., Tseng, R.J.H., and Yang, Y., Adv. Mater., 17, 1440, 2005. With permission.)
performed over a 14 h period and ran until the parameter analyzer reached its 10,000 point limit at which time no significant change in conductivity is witnessed (inset, Figure 8.24). The PANI nanofiber–gold nanoparticle device exhibits very fast response times while switching, less than 25 ns for switching on and off. Cycle tests on the device are conducted to observe their lifetime, and the parameters used for the write–read–erase voltages are as follows: the ‘‘write’’ voltage is 4.8 V, the ‘‘read’’ voltage is 1.2 V, and the ‘‘erase’’ voltage is 6 V. The current is measured at the ‘‘read’’ voltage
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10−9
Capacitance (F)
OFF-state ON-state
10−10
10−11 10
100
1000
104
105
106
Frequency
FIGURE 8.22 Typical frequency dependence of capacitance of the device Al=PS þ PCBM þ TTF=Al in high (ON) and low (OFF) conductivity states. (From Chu, C.W., Ouyang, J., Tseng, R.J.H., and Yang, Y., Adv. Mater., 17, 1440, 2005. With permission.)
FIGURE 8.23 TEM image of the polyaniline nanofiber=gold nanoparticle composite. The black dots are 1 nm gold nanoparticles contained within 30 nm diameter polyaniline nanofibers. (From Tseng, R.J.H., Huang, J., Ouyang, J., Kaner, R., and Yang, Y., Nano. Lett., 5, 1077, 2005. With permission.)
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10−3 A 10−4
B
NDR ON
C
10−6
10−4 10−5
10−7 OFF
10−8
Current (A)
Current (A)
10−5
ON state
10−6 10−7 10−8
OFF state
10−9
10−9
10−10 100
101
102
103
104
105
Time (s)
10−10 0
1
2
3 4 Voltage (V)
5
6
7
FIGURE 8.24 Current–voltage characteristics of the polyaniline nanofiber–gold nanoparticle device. The potential is scanned from (A) 0 to þ4 V, (B) þ4 to 0 V, and (C) 0 to þ4 V. Between þ3 and þ4 V, a region of negative differential resistance (NDR) is observed. The inset shows the retention time test of the ON state (top) and OFF state (bottom) currents when biased at þ1 V every 5 s. (From Tseng, R.J.H., Huang, J., Ouyang, J., Kaner, R., and Yang, Y., Nano. Lett., 5, 1077, 2005. With permission.)
(1.2 V) for the device in the high (ON) and low (OFF) conductivity states and was recorded to be 105 A and 106 to 107 A, respectively. A readily distinguishable ON=OFF ratio around 20 is always maintained. The switching mechanism in PANI nanofiber–gold nanoparticle devices is similar to the all-organic donor–acceptor system discussed in Section 8.2 and appears to be a result of an electric field-induced charge transfer effect between the PANI nanofibers and the gold nanoparticles. Under a sufficient electric field, electrons that reside on the imine nitrogen of the polyaniline may gain enough energy to surmount the interface between the nanofibers and the gold nanoparticles and transfer onto the gold nanoparticles (Figure 8.25). Consequently, the gold nanoparticles become more negatively charged whereas the polyaniline nanofibers become more positively charged. The conductivity of the polyaniline nanofiber– gold nanoparticle composite will increase dramatically after the electric field-induced charge transfer. To provide experimental observation corroborating this ideology, first, x-ray photoelectron spectra taken on the composite shows a shift from 399.2 to 399.7 eV for the N1S core electrons compared to undoped emeraldine base polyaniline indicating that the nitrogen in the PANI nanofiber–gold nanoparticle composite is partially positively charged. At the same time, the binding energy of the gold electrons (4f5=2) decreases from 87.7 to 87.5 eV, indicating that a partial negative charge resides on the gold nanoparticles. Second, our assumption of an interface between the polyaniline nanofibers and gold nanoparticles seems reasonable, since without such an interface instability of the device through rapid charge recombination would be expected. Additionally, since our device exhibits NDR, a mechanism involving filament formation is unlikely as discussed by Scott and coworkers [16]. As miniaturization trends continue, and memory devices continually reach higher densities, this material demonstrates unique appeal due to the inherent nanoscale dimensions of the individual components of the nanocomposite. In order to show that even extremely small amounts of the nanocomposite material display the memory effect, a nanoscale writing or reading process is carried out using a conductive AFM tip in direct contact with the polyaniline nanofiber–gold nanoparticle thin
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Auδ− e– δ+ N
H N
N
N H
n
Emeraldine base
FIGURE 8.25 Schematic structure of a polyaniline nanofiber–gold nanoparticle after application of þ3 V. An increase in charge transfer from polyaniline to the gold nanoparticles is believed to occur. (From Tseng, R.J.H., Huang, J., Ouyang, J., Kaner, R., and Yang, Y., Nano. Lett., 5, 1077, 2005. With permission.)
film (in this experiment no polyvinyl alcohol is used). An external electric voltage is applied between the AFM tip, which is acting as the top electrode, and the bottom Al electrode to examine if the electrical bistable switching effect is present. Initially, the surface of the film is scanned to obtain an image of the surface morphology of the composite (Figure 8.26, lower left) and from this initial scan a ‘‘bump’’ containing a group of nanofibers is located (upper right, Figure 8.26). The AFM tip is situated directly over the top of the bump and a voltage scan from 0 to 5 V is applied, while the current is measured.
TUNA current pA
100
50
0
−4
500
−3 −2 −1 DC sample bias (V)
400 300 200 100
0
100
200
300
400
500
nm
FIGURE 8.26 (See color insert following page 8-22.) Conductive atomic force microscopy of the polyaniline nanofiber–gold nanoparticle composite. A conductive atomic force microscope tip is used to first perform the morphology scan on the polyaniline nanofiber–gold nanoparticle composite film and then to carry out the electrical characterization. The AFM tip is parked on the top of the polymer bump and a voltage scan is taken from 0 to 5 V, while current is measured. The electrical bistability of the polymer composite film using the nanoscale tip is evident. (From Tseng, R.J.H., Huang, J., Ouyang, J., Kaner, R., and Yang, Y., Nano. Lett., 5, 1077, 2005. With permission.)
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Electrical bistability is indeed observed in the nanocomposite film verifying that the nonvolatile memory effect is present even at nanoscale dimensions. This observation will continue to drive researchers towards developing future organic memory devices.
8.5
Conclusion
This paper provides an introduction and review of electrically addressable bistable devices based on polymer films. Several different material and device structures are shown, including polymers blended with synthesized metal nanoparticles, composites of polymer nanofibers decorated with nanoparticles, and donor–acceptor complexes. Conductance switching is observable from the I–V characteristics, which show a conductance change of at least three orders of magnitude. Transitions between the ON and OFF states occur in the nanosecond regime, and such fast response times suggest that the switching is due to electric field-induced charge transfer. This effect is discussed for each system presented in this chapter, and details of retention time and write–read–erase cycles for memory performance are also provided.
Acknowledgment The authors would like to thank Prof. Richard Kaner of the Chemistry Department at UCLA for his collaboration and for providing the PANI nanofiber. Technical discussions with Dr. Chuck Szmanda of Rohm & Haas, Prof. Qibing Pei and Prof. Fred Wudl of UCLA, and Dr. Ed Chandross of Bell Laboratories are also acknowledged. Finally, we would like to express our ultimate thanks to the financial support for the past several years on the polymeric memory devices. The Air Force Office of Scientific Research provides the funding for the polymer memory devices. The PANI=Au NP memory device is supported by the UCLA FENA-MACRO Center with funds from the Semiconductor Research Corporation (SRC) and the Defense Advanced Research Project Agency (DARPA). The interface study of this work is sponsored by the National Science Foundation.
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19. 20. 21. 22. 23. 24. 25. 26. 27. 28. 29. 30. 31. 32. 33. 34. 35. 36.
37. 38. 39. 40. 41. 42. 43. 44. 45. 46. 47. 48. 49. 50. 51. 52. 53. 54. 55. 56. 57. 58. 59. 60. 61.
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Mahapatro, A.K., R. Agrawai, and S. Ghosh. 2004. J Appl Phys 96:3583. Tondelier, D., K. Lmimouni, D. Vuillaume, C. Fery, and G. Haas. 2004. Appl Phys Lett 85:5763. Chen, J., and D. Ma. 2005. Appl Phys Lett 87:023505. Mo¨ller, S., C. Perlov, W. Jackson, C. Taussig, and S.F. Forrest. 2003. Nature 426:166. Furukawa, T. 1997. Adv Colloid Interf Sci 71–72:183. Ouyang, J., C.W. Chu, D. Sievers, and Y. Yang. 2005. Appl Phys Lett 86:123507. Ouyang, J., C.W. Chu, R.J.H. Tseng, A. Prakash, and Y. Yang. 2005. Proc IEEE 93:1287. Chu, C.W., J. Ouyang, R.J.H. Tseng, and Y. Yang. 2005. Adv Mater 17:1440. Tseng, R.J.H., J. Huang, J. Ouyang, R. Kaner, and Y. Yang. 2005. Nano Lett 5:1077. Hostetler, M.J., J.E. Wingate, C.J. Zhong, J.E. Harris, R.W. Vachet, M.R. Clark, J.D. Londono, S.J. Green, J.J. Stokes, G.D. Wignall, G.L. Glish, M.D. Porter, N.D. Evans, and R.W. Murray. 1998. Langmuir 14:17. Chen, Y., D.A.A. Ohlberg, X.M. Li, D.R. Stewart, R.S. Williams, J.O. Jeppesen, K.A. Nielsen, J.F. Stoddart, D.L. Olynick, and E. Anderson. 2003. Appl Phys Lett 82:1610. Tsujioka, T., and H. Kondo. 2003. Appl Phys Lett 83:937. Henish, H.K., and W.R. Smith. 1974. Appl Phys Lett 24:589. Segui, Y., B. Ai, and H. Carchano. 1976. J Appl Phys 47:140. Wang, W., T. Lee, and M.A. Reed. 2003. Phys Rev B 68:035416. Prout, C.K., and A.G. Wheeler. 1967. J Chem Soc A 469. Castellano, E., and C.K. Prout. 1971. J Chem Soc A 550. Adams, D.M., L. Brus, C.E.D. Chidsey, S. Creager, C. Creutz, C.R. Kagan, P.V. Kamat, M. Lieberman, S. Lindsay, R.A. Marcus, R.M. Metzger, M.E. Michel-Beyerle, J.R. Miller, M.D. Newton, D.R. Rolison, O. Sankey, K.S. Schanze, J. Yardley, and X.Y. Zhu. 2003. J Phys Chem B 107:6668. Ipe, B.I., K.G. Thomas, S. Barazzouk, S. Hotchandani, and P.V. Kamat. 2002. J Phys Chem B 106:18. Oyamada, T., H. Tanaka, K. Matsushige, H. Sasabe, and C. Adachi. 2003. Appl Phys Lett 83:1252. Mo, X.-L., G.-R. Chen, Q.-J. Cai, Z.-Y. Fan, H.-H. Xu, Y. Yao, J. Yang, H.-H. Gu, and Z.-Y. Hua. 2003. Thin Solid Films 436:259. Chen, S., R.S. Ingram, M.J. Hostetler, J.J. Pietron, R.W. Murray, T.G. Schaaff, J.T. Khoury, M.M. Alvarez, and R.L. Whetten. 1998. Science 280:2098. Hicks, J.F., A.C. Templeton, S.W. Chen, K.M. Sheran, R. Jasti, R.W. Murray, J. Debord, T.G. Schaaf, and R.L. Whetten. 1999. Anal Chem 71:3703. Frisch, M.J., et al. 2003. Gaussian 03, Revision B.05. Pittsburgh, PA: Gaussian, Inc. Kubo, R. 1962. J Phys Soc Jpn 17:975. Chen, S., R.W. Murray, and S.W. Feldberg. 1998. J Phys Chem B 102:9898. Bru¨tting, W., S. Berleb, and A.G. Mu¨ckl. 2001. Org Electron 2:1. Rhoderick, E.H., and R.H. Williams. 1988. Metal–semiconductor contacts. Oxford: Clarendon Press. Laurent, C., and E. Kay. 1988. J Appl Phys 64:336. Vollmann, W., and H.U. Poll. 1975. Thin Solid Films 26:201. Martı´n, N., E. Ortı´, L. Sa´nchez, P.M. Viruela, and R. Viruela. 1999. Eur J Org Chem 1239. Brbec, C.J., N.S. Sariciftci, and J.C. Hummelen. 2001. Adv Funct Mater 15:11. Bryce, M.R. 1999. Adv Mater 11:11. Ferrais, J., D.O. Cowan, V. Walatka, and J.H. Perlstein. 1972. J Am Chem Soc 95:948. Zheng, L., Q. Zhou, X. Deng, M. Yuan, G. Yu, and Y. Cao. 2004. J Phys Chem B 108:11921. Huang, J., S. Virji, B.H. Weiller, and R.B. Kaner. 2003. J Am Chem Soc 25:314. Huang, J., and R.B. Kaner. 2004. J Am Chem Soc 126:851. Virji, S., J. Huang, R.B. Kaner, and B.H. Weiller. 2004. Nano Lett 4:491. Huang, J., S. Virji, B.H. Weiller, and R.B. Kaner. 2004. Chem Eur J 10:1314. Wang, J., K.G. Neoh, and E.T. Kang. 2001. J Colloid Interf Sci 239:78. Smith, J.A., M. Josowicz, and J. Janata. 2003. J Electrochem Soc 150:E384. Le, J.D., Y. He, T.R. Hoye, C.C. Mead, and R.A. Kiehl. 2003. Appl Phys Lett 83:5518. Chen, J., J. Su, W. Wang, and M.A. Reed. 2003. Physica E 16:17.
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9 Electroactive Polymers for Batteries and Supercapacitors 9.1
Introduction......................................................................... 9-1 Terminology . Differences between Batteries and Supercapacitors . Ion Movement
9.2
Electroactive Polymer-Based Batteries............................... 9-5 Applications . General Description of Batteries . Review of Electroactive Polymer Battery Literature . Summary of Electroactive Polymer Battery Research
9.3
Jennifer A. Irvin, David J. Irvin, and John D. Stenger-Smith*
9.1
Electroactive Polymer-Based Supercapacitors................. 9-10 Applications . General Description of Supercapacitors . Supercapacitor Definitions . Performance Implications of Supercapacitor Device Types . Hypothetical Supercapacitor Outputs . Review of Supercapacitor Literature . Critical Analysis of Supercapacitor Research
9.4
Conclusions........................................................................ 9-20
Introduction
Current charge storage technologies cannot meet the growing power requirements of a vast range of new and improved electronic devices. Researchers have therefore turned to alternative technologies, such as those based on electrically active polymers, to fulfill these needs. The oxidation and reduction (redox) processes in electroactive polymers (EAPs) make it possible to use these polymer materials as charge storage devices, either as battery electrodes or as supercapacitors. The potential for reduced cost, weight, and environmental impact of EAP electrodes relative to the metals and metal oxides that are traditionally used in such devices makes these polymers attractive alternatives. While inorganic options are limited, EAPs can be tailored to provide specific properties, such as conductivity, voltage window, storage capacity, porosity, reversibility, and chemical and environmental stability. This chapter will present the fundamentals of EAP-based charge storage devices, namely secondary batteries and supercapacitors. Design considerations and characteristic performance parameters for these devices will be discussed, and research conducted in this area will be reviewed.
9.1.1 Terminology Before embarking on a review of EAPs for batteries and supercapacitors, it is necessary to define, discuss, and describe several terms for the sake of consistency. While most of this terminology will be familiar to those in the EAP field, there are several terms that are used interchangeably or that differ for batteries and supercapacitors, and clarification is needed. *The views presented are those of the authors and do not necessarily represent the views of DoD or its components.
9-1
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The terminology used to describe redox processes in n-doping polymers is often ambiguous. For clarity in this chapter, oxidation of polymers is defined as the removal of electrons from the polymer in its neutral (uncharged or undoped) state. This oxidized (p-doped) polymer is positively charged and most commonly referred to as the true conducting version of an EAP. When this oxidized polymer is converted back to its neutral form by the addition of electrons (often referred to as reduction of the oxidized polymer), it will be referred to as neutralization of the oxidized polymer, rather than reduction. Reduction of EAPs is defined as the addition of electrons to the neutral polymer. This reduced (n-doped) polymer is negatively charged. When this reduced polymer is converted back to its neutral form by the removal of electrons (often referred to as oxidation of the reduced polymer), it will be referred to as neutralization of the reduced polymer to avoid confusion. The classical definition of a battery is a combination of two or more galvanic cells electrically connected to work together to produce electric energy. However, today a battery can be defined as any of a family of electrochemical charge storage devices from a lead-acid car battery to hybrid electrostatic and electrochemical single cells. Primary batteries are assembled in the charged state, used until fully discharged, and discarded. Secondary batteries are rechargeable by reversing the current flow, so that they are typically assembled in the discharged state, charged before use, and then repeatedly discharged and recharged. Oxidation occurs at the negative electrode, or anode, of the cell whereas reduction occurs at the positive electrode, or cathode. While terminology for batteries is relatively straightforward, capacitor terminology is not, because of several types of capacitors and nomenclature that is often used interchangeably among types. Doublelayer capacitor, electrochemical capacitor, redox capacitor, supercapacitor, and ultracapacitor are all used to describe a variety of charge storage devices; these terms differentiate this class of materials from traditional electrolytic and electrostatic capacitors. Electrostatic effects are responsible for charge storage in double-layer capacitors: the stored electrical energy is based on the separation of charged species in the electrical double layer between an electrolyte and an electron conductor (commonly a carbonaceous material). Extensive research in this area has been reviewed by several groups [1–4] and will not be the focus of this chapter. Instead, this chapter will focus on redox capacitors, also known as electrochemical capacitors, devices in which charge is stored chemically via oxidation and reduction processes in the active materials. Much work has been done on metal oxide-based redox capacitors, particularly ruthenium oxide, but this will not be discussed in this chapter; several reviews have been published in this area [1–3,5,6]. Instead, EAP-based redox capacitors will be the focus of this chapter. For another review on EAPs for charge storage, see Ref. [7]. The terms supercapacitor and ultracapacitor are used to describe any double layer or redox capacitor with specific energy and specific power intermediate to batteries and traditional capacitors. Typically, ultracapacitor refers to a device comprised of two carbonaceous electrodes whereas supercapacitor refers to a similar device in which the two carbonaceous electrodes are catalyzed with metal oxides such as RuO2. This chapter will use the term supercapacitor to describe EAP-based capacitors, since that seems to be the most commonly used term for such materials. Another charge storage configuration uses an EAP electrode and a battery-type carbonaceous electrode in what is known as a hybrid device (however, outside of the EAP-based supercapacitor field, hybrid may refer to the combination of a battery electrode such as nickel hydroxide with a carbon electrode) [1]. Additional charge storage terminology is needed to characterize the effectiveness of the hybrid devices. The energy content of a system is described by specific energy (Watt-hour per Kilogram, Wh=kg, also known as gravimetric energy density) and energy density (Wh=L, also known as volumetric energy density) whereas the discharge rate of a system is described by its specific power (W=kg, also known as gravimetric power density) or its power density (W=L, also known as volumetric power density). In Ragone plots, specific power is plotted against specific energy; the differences between batteries, supercapacitors, and capacitors are readily visualized in these plots (Figure 9.1) [1]. The most important metric for all types of capacitors is charge storage ability. This is described in terms of capacity (coulombs per gram) or capacitance (capacity per volt, F=g or F=cm3). Strictly speaking, capacitance and specific capacitance refer to charge per volt accumulated in the double layer
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107 Capacitors Specific power (W/kg)
106 Combustion/ turbine
105 104 103
Supercapacitors
102
Batteries
10 1 0.01
0.1
Fuel cells
1 10 100 Specific energy (Wh/kg)
1000
FIGURE 9.1 Ragone plot for different types of power sources. (Adapted from M. Winter and R.J. Brodd, Chem. Rev., 104, 4245, 2004. Copyright 2004, American Chemical Society. With permission.)
whereas pseudocapacitance refers to charge storage per volt in association with a redox reaction. For a thorough treatment of pseudocapacitance, see the discussion by Conway et al. [8]. Unfortunately, many researchers use the terms capacity, capacitance, and pseudocapacitance interchangeably. Capacitance values are obtained from cyclic voltammetry or from impedance spectroscopy, although one report [9] noted that values obtained using impedance spectroscopy are smaller than those obtained from cyclic voltammetry. The authors attribute this difference to conformational changes in polymers during oxidation and neutralization. Coulombic efficiency and voltage efficiency are usually expressed in terms of percentage of total charge or voltage available. Cycle life is a loosely defined parameter that usually specifies a number of cycles before a certain percentage loss in capacity. It is also useful to consider charge density, which is the mass of active polymer required per unit charge (Ah=kg). Similarly, total device charge density includes the mass of active polymer, electrolyte, solvent, charge collection and distribution electrodes, and any solid supports or encapsulation materials. Charge dissipation, or self-discharge, rate is usually characterized by a percentage loss in capacity per unit time, e.g., 1% loss per day. Other terms used to characterize charge storage devices include open circuit voltage (OCV, the voltage across the device when no external current flows) and internal resistance or impedance (equivalent series resistence (ESR), the resistance or impedance the device exhibits to current flow).
9.1.2 Differences between Batteries and Supercapacitors Figure 9.2 [10] is a useful guide illustrating the distinctions among batteries, capacitors, and supercapacitors. Batteries can hold large quantities of energy, but are not capable of delivering energy at very high rates (high power). Capacitors can deliver energy at very high rates (very high power), but are not capable of holding large quantities of energy. Supercapacitors are intermediate in power and energy delivery. The difference between EAP-based batteries and EAP-based supercapacitors is not clear. According to Kotz and Carlen [3], in terms of voltage transient during charge and discharge, and with respect to the cyclic voltammetry, supercapacitors should be considered batteries rather than capacitors. However, in comparison with metal oxides, they are capacitors with one redox peak rather than the rectangular cyclic voltammograms typical of metal oxides. Conway et al. [8] suggest that the utilization of pseudocapacitance in supercapacitors makes them a transition case between double-layer electrostatic charge storage (traditional capacitors) and electrochemical battery cell charge storage since pseudocapacitance takes place through Faradaic processes (typical of batteries) accompanied by a (usually much smaller) nonFaradaic double-layer capacitance (typical of traditional capacitors) occurring at the interface between the polymer and the electrolyte.
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1.2
Depth of discharge
1 0.8 0.6
Battery Supercapacitor Capacitor
0.4 0.2 0 0.001
0.01
0.1
1
10
100
1000
Frequency (Hz)
FIGURE 9.2 A comparison of discharge rates in batteries, capacitors, and supercapacitors. (Adapted from M. McKubre, An evaluation of supercapacitor technology, Office of Naval Research Grand Challenges Workshop, SRI, 1999. With permission.)
Supercapacitors are similar to batteries in the following ways: . . .
Both contain or require a cathode, an anode, electrodes, and an electrolyte. Both have complementary oxidation and reduction reactions occurring at each electrode. Both store more energy than capacitors.
Supercapacitors are similar to capacitors in the following ways: . . .
Both are capable of delivering more power than batteries. Both deliver less energy than batteries. During capacitor and supercapacitor charging and discharging, only the ions move in and out of a thin layer.
The similarities between the two types of devices are so strong that devices described in the literature as one type may later be referred to as the other. For instance, Arbizzani et al. [11] reported the preparation of polydithieno[3,4-b:30 ,40 -d]thiophene devices, described as supercapacitors. Two years later, Novak et al. [12] included those devices in a review on EAP-based batteries, along with a dozen other devices that might be described as either batteries or supercapacitors. An early report of a polypyrrole device constructed by Mohammadi et al. [13] was originally described as a secondary battery, but it could arguably be classified as a supercapacitor. For this chapter, we have chosen to stay consistent with the terminology chosen by the original researchers.
9.1.3
Ion Movement
It is well known that changes in redox states of EAPs require movement of ions in order to maintain electroneutrality. Electrochemical neutralization of p-doped polymers can be anion-dominant (anions are expelled from the polymer film), cation-dominant (cations diffuse into the polymer film), or a combination of both anion and cation movements [14,15]. The dominant process varies from polymer to polymer and is strongly affected by ion and solvent choices as well as by electropolymerization conditions and film thickness [16]. Typically for p-doping polymers, anion transport dominates when the anion is small and highly mobile, but cation transport dominates when the anion is large and immobile [17–19]. For instance, redox processes in polypyrrole are anion-dominant in most cases, but when poly(styrenesulfonate) is chosen as the counterion, cation transport is the dominant process [20–23]. Ion transport in n-doping polymers needs further investigation, but it seems likely that cation transport dominates when the cation is small and highly mobile whereas anion transport dominates
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when the cation is large and immobile. Ion transport processes are studied using a variety of techniques, including electrochemical quartz crystal microbalance (EQCM) [16,17,24–27], rotating ring-disk electrode voltammetry [28–30], impedance techniques [15], luminescence probes [31], and spectroscopic techniques [32,33]. Cation transport is dominant for several ionic liquid electrolytes, even though the cations are larger than the anions; the reasons for this are unknown [34,35]. Differences in cations have been shown to result in changes in morphology, electrochemical stability, and oxidation and reduction potentials [36]. A report by Stenger-Smith et al. [37] indicates that the cation can also play an important role in the cycle lifetime and cycle speed of polymer supercapacitors.
9.2
Electroactive Polymer-Based Batteries
Batteries using EAPs for the anode or the cathode have received considerable attention over the past 25 years. The materials show mixed electronic and ionic conductivity during charge and discharge, and they can exhibit high conductivities, flexible morphologies, chemical stability, ease of manufacturing, and low cost. Several reviews on EAPs for charge storage have been published in the last decade [8,7,12]. While a wide range of polymeric materials is used in the study and manufacture of batteries, we are restricting the scope of this work to include only conjugated, delocalized EAPs as the anode or the cathode. Novak et al. [12] detailed a comprehensive review covering EAPs, carbons, and redox materials outside the scope of this work.
9.2.1 Applications Portable consumer electronics: The drive toward portable electronic devices has led to an increase in the demand and market share of rechargeable batteries. Batteries can power everything from cars to hearing aids. Li-ion and Ni-hydride cells are approaching their practical limits. To obtain higher power densities, new materials and charge storage systems are required. Military applications: Until recently, rechargeable batteries were not reliable enough for military use, and the costs and logistics of recharging were prohibitive. Now, however, the average ground force soldier requires approximately 150 lb of batteries for a 2-week mission, so the military is becoming increasingly dependant on rechargeable battery-powered devices. Much of this increased demand stems from an increase in the number of portable electronic military devices, such as GPS locators and night vision goggles. Electric vehicles: While there is a strong drive to produce electric cars, the internal combustion engine still has higher power and energy densities as well as increased cruising range (Figure 9.1). Consumers want more power, faster recharge times, and lower costs. The current systems, such as lead-acid and Nihydride batteries, have been engineered to their practical limits. Only new materials for batteries can meet all of the demands from current and future consumers. In applications where weight is at a premium, such as consumer vehicles, aircraft (both manned and unmanned), and missiles, high charge density, low physical density batteries (such as polymer-based systems) will be needed.
9.2.2 General Description of Batteries Primary batteries, by definition, are produced in their charged form and can only release charge. Secondary batteries, by contrast, can be charged and discharged more than once. Within these two broad classes of batteries are several subclasses: lead-acid, aqueous, lithium ion, solid polymer, etc. Regardless of the type of battery, the essential elements are the same (Figure 9.3): an electrically conductive anode, an ion-conductive electrolyte, an optional physical separator (to prevent short circuits) within the electrolyte layer, and an electrically conductive cathode. In EAP-based batteries, a p-doping polymer is used as the anode, and an n-doping polymer is used as the cathode. To store electrical energy, ions must be moved from the electrolyte into the electrode to help balance charge, or conversely be expelled if the ion is of the same charge as those created in the electrode. When
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e−
Charge collectors Anode Separator Electrolyte Cathode
FIGURE 9.3 Representation of the generic construction of a battery or supercapacitor. Electroactive polymers could be used as anode or cathode.
this occurs, there is substantial deformation of the solid structure of the electrode (especially when the materials, such as lead, dissolve and redeposit). In rigid systems, the deformation leads to mechanical and then electrical failure. The inherent plasticity of polymers minimizes this effect, thus making polymers an excellent material for secondary battery electrodes. A majority of the polymers currently used in batteries are used as electrolytes, such as polyethylene oxide in lithium–polymer batteries, but serving only as ion conductors they are outside the scope of this work. Reviews on ion-conducting polymers can be found in the literature [38–41]. For batteries, the power density is lower than the energy density because the discharge process requires the movement of ions, and thus mass. Ion movement is often accompanied by a phase change, 0 such as the deposition of lead metal (Pb2þ (aq) ¼ Pb(s) ). This movement of mass slows the process, decreasing the effective power (i.e., energy=unit time). The charge or discharge rate is directly related to the rate at which ions can be incorporated and expelled from the electrode material, barring electrostatic charging. Secondary batteries have demonstrated specific energies of 10–500 Wh=kg with power densities of 10–1000 W=kg, although power densities for polymer electrode-based batteries are rarely reported over 500 W=kg. When evaluating new materials in the literature, a few key measurements or estimates are most relevant. First is the discharge voltage of the cell. Values between 2 and 4 V are obtainable and important for the power density. It is also important to note whether the data are from a two-electrode device or a three-electrode test cell; the two-electrode device more closely emulates the final use. Next is the charge density; values greater than 200 Ah=kg of polymer are good targets. Another useful value is the total device charge density, which includes all device components; this value is typically at least an order of magnitude lower than the charge density based on polymer weight. The coulombic and voltage efficiency of the device are also important, with values of >85% as useful minimum efficiencies for each. Some device-related parameters include number of charge–discharge cycles (cycle life) and charge dissipation rate (self-discharge rate). Li-ion batteries are the current state of the art with 1000 cycles, therefore this should be a threshold value. Since secondary batteries are rechargeable, a self-discharge rate of approximately 1% or less per day should be acceptable. Another
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important factor in comparing secondary batteries is the relationship between depth of discharge (DOD) and expected life cycle. The accessible capacity in secondary batteries may be considerably lower than the theoretical capacity. Some material-dependent amount of the battery’s capacity must be reserved to ensure reversibility, thus lowering the DOD [42,43].
9.2.3 Review of Electroactive Polymer Battery Literature Since the rule of thumb for the practical energy content of a secondary battery is 25% of its theoretical value [43], the usable capacity and the energy density need to be increased to become practical in highuse arenas such as electric vehicles. New materials, electrodes, cell designs, and charge collection schemes will be required for the next generation of high power- and energy-density batteries. EAPs investigated for use in batteries are for the most part the same as those used in supercapacitors. The structures of these polymers are given in Figure 9.5. 9.2.3.1 Polyacetylene Polyacetylene (PA, Figure 9.4A) is probably the most studied form of the conductive polymers, and it has been used in a variety of battery types. PA can be used as the anode or cathode since it both n-dopes and p-dopes [44–49]. For the cathode, the predominant counterions used are ClO4 [44,50,51] and I [45] with BF4 [52], PF6 [53], and AsF6 [47,54] used to a lesser extent. Electrochemically doping the neutral polymer in an electrolyte solution, typically in propylene carbonate (PC) or tetrahydrofuran (THF), usually results in doping levels between 0.05 and 0.10 anions per repeat unit. As a result, test cells can reliably obtain specific charge densities of 100–300 Ah=kg. The specific energy for PA-based electrodes ranges between 100 and 300 Wh=kg. The OCV of a cell using PA as the cathode and Li as the anode is between 3.5 and 3.9 V [44,55,56]. When an all PA cell is used and Liþ is used to dope the anode, the OCV is stable at ca. 3.5 V [46,48,57], but drops significantly to ca. 2.5 V if tetrabutylammonium (Bu4Nþ) is used as the counterion [44,53,57]. While extensive research was conducted on polyacetylene batteries in the 1980s and 1990s [58–61], little has been done recently. This is likely due to stability issues and processing difficulties. Since a majority of PA is synthesized via Ziegler–Natta-type catalysts [62], trace amounts of Ti and Al, as well as other metals, are always present. These metals may decrease long-term stability or cause other side reactions if they migrate and react. PA is typically formed as fibrils or powders, which are insoluble. The as-cast material must then be electrochemically doped to produce working electrodes. The inaccessible (e.g., crystalline) portions of the polymer may account for the lower doping levels when compared with electrochemically grown polymers. 9.2.3.2 Polyaniline Polyaniline (PANI, Figure 9.4B) has also been extensively studied for use in batteries. With few exceptions [63–65], PANI is used as the cathode. PANI works well in aqueous and nonaqueous cells, although the nonaqueous cells have a much higher OCV. For aqueous cells, a Zn anode will yield an OCV between 1.0 and 1.5 V [63,66–69] whereas a nonaqueous cell with a Li anode will yield an OCV of 3.0–4.0 V [70–85]. In addition to Zn and Li, magnesium has also been used as an anode in PANI-based cells [86,87]. It should be noted that in the aqueous systems, acids are commonly used as electrolytes, whereas in Li-based cells, ClO4 is the predominant counterion. There has also been a recent surge in composite electrodes that capitalize on the best properties of the individual components, thus expanding the voltage range and reliability [78,88–93]. The charge density for PANI-based cells reliably ranges from 50–150 Ah=kg, with energy densities of 100–350 Wh=kg. With the electrochemically grown films, there are no major barriers for using PANI in batteries, although chemically synthesized PANI has minor solubility and processability issues. 9.2.3.3 Polypyrrole Battery electrodes-based on polypyrrole (PPy, Figure 9.4C) are similar to electrochemically grown PANI [94–98]. They exhibit OCVs ranging from 3.0 to 4.0 V in nonaqueous Li-anode cells with charge
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R
n (Q)
FIGURE 9.4 Polymers investigated for use in charge storage devices include: (A) polyacetylene (PA), (B) polyaniline (PANI), (C) polypyrrole (PPy), (D) polythiophene (PT), (E) poly-3-methylthiophene (PMT), (F) poly-3phenylthiophene (PPT) and its derivatives (PFPT, etc.), (G) polydithieno[3,2-b:20 ,30 -d]thiophene, (H) poly-pphenylene (PPP), (I) polydiaminoanthraquinone (PDAAQ), (J) poly-3,4-ethylenedioxythiophene (PEDOT), (K) poly{1,3-bis[20 -(30 ,40 -ethylenedioxy)thienyl]-benzo[c]thiophene-N-2(-ethylhexyl-4,5-dicarboximide} (PDEI), (L) poly-3,4-propylenedioxythiophene (PProDOT), (M) polydithieno[3,4-b:30 ,40 -d]thiophene (PDTT), (N) poly-3-(ptrimethylammoniumphenyl)bithiophene, (O) polycyclopenta[2,1-b;3,4-b0 ]dithiophen-4-one, (P) poly(4,40 -bicyclopenta[2,1-b;3,4-b0 ]dithiophenylidene), and (Q) poly(cyano-substituted diheteroareneethylenes).
densities between 50 and 170 Ah=kg and energy densities of 50–350 Wh=kg. As in PANI, ClO4 is the dominant counterion in PPy-based cells [99–107], although BF4 [102,108,109] has also been used effectively. The theme of composite materials is continued in PPys with modest improvements in properties [110–114]. 9.2.3.4 Polythiophene and Its Derivatives Polythiophene (PT, Figure 9.4D) is the most easily functionalized of the polymer systems surveyed here. While the bulk of polythiophenes have been used as cathode materials, there are several systems that have been found to n-dope; these are used as anode materials. In general, PTs typically exhibit a specific charge of 25–100 Ah=kg and a specific energy of 50–325 Wh=kg. Polythiophene electrochemically synthesized from bithiophene [115–117] or terthiophene [118] exhibits, as predicted, cleaner electrochemistry and more stable battery materials. Poly(3-methylthiophene) (PMT, Figure 9.4E) is well studied [119–122], with specific charge ca. 90 Ah=kg and one report of specific energy of 326 Wh=kg [123]. In nonaqueous
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systems, the OCV ranges from 3.2 to 4.2 V [95]. The fused ring system of dithienylthiophene has also been well studied [11,124–128] with an OCV of ca. 3.3 V [124]. Alkoxy-substituted polythiophenes [129], poly(3-phenylthiophene) (PPT, Figure 9.4F, where X ¼ H) [130], and poly(dithienylethylenes) [131] have been incorporated into cathodes. Once again, composite electrodes are used to improve the good properties of PT-based electrodes [112,132–134]. As for anode materials, three of the better materials are PPT, poly(3-(4-fluorophenyl)thiophene) (PFPT, Figure 9.4F, where X ¼ F), and polydithieno[3,2-b:20 ,30 -d]thiophene (Figure 9.4G). All have been shown to work as anodes and cathodes in different solvent and electrolyte systems. For PPT, tetraethylammonium hexafluorophosphate (Et4NPF6) in acetonitrile (AN) yielded a device with a specific energy of 25 Wh=kg [130]. In contrast, a device with PFPT as both anode and cathode using tetramethylammonium triflate (Me4NCF3SO3) in AN is reported to have a lower specific energy (18 Wh=L) but a very high specific power of 35,000 W=kg [135]. 9.2.3.5 Poly-p-phenylenes While poly-p-phenylene (PPP, Figure 9.4H) is not as highly redox active as the previous systems, it works well in alkali metal batteries, partly due to its crystallinity and resemblance to graphite [136,137]. PPP can be used as a cathode with lithium [138,139]. Using PPP as the cathode and Li as the anode, OCV of between 3.2 and 4.5 V are reported [140]. For all PPP=Li batteries, the specific charge is between 20 and 140 Ah=kg but can be driven higher (>150 Ah=kg) with the addition of carbon [141]. The specific energy is reported as high as 300 Wh=kg [138]. A series of counterions were studied for PPP=Li batteries including BF4, ClO4, PF6, AsF6, and triflate [139], resulting in PF6 and AsF6, performing better than the others due to their size and ‘‘soft’’ nature, although these observations may be skewed since the polymers were electropolymerized in LiAsF6. 9.2.3.6 Electrolytes None of the commonly used electrolytes are universally satisfactory for use in nonaqueous batteries. While water-based batteries have been demonstrated, the OCVs of these cells are dramatically lower than those for nonaqueous systems. Most nonaqueous solvent and electrolyte systems possess electrochemical stability of ca. 4 V, and the long-term stability and reactivity to the delocalized anions and cations on the polymers in their charged (doped) state needs further study. In addition, the use of volatile organic solvents is becoming more regulated due to health as well as fire hazards. For room temperature and above, moving to solid electrolyte systems may be a viable approach, but as the temperature of these cells goes down, so does the ion mobility and therefore the output. For low-temperature applications, new polyelectrolytes with improved ion mobilities or ionic liquids may be needed. For a thorough review of battery electrolytes, see Xu [142].
9.2.4 Summary of Electroactive Polymer Battery Research A number of the current cathode materials show promise based on coulombic efficiency, cycling stability, and voltage window. Coupled with lower density and fast charging or discharging rates, these materials will likely be further commercialized. Since PA, PANI, PPy, and PT derivatives all perform well, the choice between them comes down to processing method, robustness of packaging, intellectual property, and cost. The critical values of cost, availability, source, voltage window, cycle life, etc., must be balanced based on the application in order to determine which polymer is best suited. As power requirements become increasingly more diverse, a variety of polymers with different discharge profiles may find specific uses in markets. The lack of robust n-doping polymers is still the largest hurdle for the realization of all-polymer batteries. While PA does n-dope and cycle, its coulombic efficiency is very low, as is its environmental stability. A few polythiophene derivatives show promise, but more research is needed. The usable capacity and energy density of EAP-based batteries need to be increased. New materials, electrodes, cell designs, and charge collections schemes will be required for the next generation of high-power and energy-density batteries.
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9.3
Conjugated Polymers: Processing and Applications
Electroactive Polymer-Based Supercapacitors
The active materials in supercapacitors can be either inorganic oxides (for example, hydrous ruthenium oxide) or EAPs. Because the charge storage mechanism in the inorganic oxides involves only the surfaces of the oxide particles, capacity is surface area-limited. EAP-based supercapacitors, on the other hand, can be designed to ensure that the entire volume of the polymer is involved in the charge storage process, and capacities could be much higher than in the oxide devices. Other advantages of the polymeric supercapacitors include lower cost, high conductivities, flexible morphologies, chemical stability, and ease of manufacturing. An EAP-based supercapacitor is generally comprised of charge collectors, electrolyte, a separator, and two polymer electrodes (Figure 9.3), which can be either the same polymer for both electrodes or different polymers for each electrode. For previous reviews on this topic, see Rudge et al. [143], Arbizzani et al. [7,144], and Shukla et al. [145].
9.3.1
Applications
With power and energy densities intermediate to those of batteries and capacitors, supercapacitors are a bridging technology that may find use in both battery and capacitor applications. Applications particularly targeted for supercapacitor use include burst power for military systems, backup power for computers, electronic fuses, and burst power for electric vehicles.
9.3.2
General Description of Supercapacitors
EAP-based supercapacitors have been divided into three types by Rudge et al. [143]. A Type I supercapacitor utilizes the same p-doping polymer for both electrodes. A Type II supercapacitor uses different p-doping polymers on each electrode. These are frequently referred to as symmetric and asymmetric supercapacitors, respectively. A Type III supercapacitor is a symmetric device utilizing the same material as both the n-doping polymer and the p-doping polymer. However, since Type III supercapacitor is the only term for a supercapacitor that incorporates an n-doping polymer, the term has since been used in the literature to describe asymmetric devices utilizing one polymer for p-doping and a different polymer for n-doping [146]. For consistency with the original definitions and for improved clarity, we propose that Type III refers specifically to a symmetric device utilizing the same polymer as both the n-doping polymer and the p-doping polymer. Asymmetric devices using one polymer for the n-doping electrode and a different polymer for the p-doping electrode would then be classified as Type IV supercapacitors. Operational differences between symmetric and asymmetric supercapacitors necessitate the need for clarity. For example, in symmetric capacitors (Type I and Type III), there is no inherent polarity. Asymmetric (Type II and Type IV) capacitors, on the other hand, are polar devices, like batteries, so that care must be taken to connect the devices in the proper orientation.
9.3.3
Supercapacitor Definitions
The voltage of a supercapacitor, defined as the potential difference between the cathode and the anode [4], can be represented using the following equation: V ¼ [Eox (anode) Ered (cathode)] þ overvoltage
(9:1)
where Eox (anode) is the oxidation potential of the p-doping polymer on the anode, and Ered (cathode) is either the oxidation potential of the p-doping polymer on the cathode (in the case of Type I and Type II supercapacitors) or the reduction potential of the n-doping polymer (in the case of Type III and Type IV supercapacitors). A practical measurement for electrochemical supercapacitors, capacitance (C) is
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defined as the charge storage capacity (Q) divided by the voltage [4]. The actual energy stored in a supercapacitor therefore simplifies to the following term: 1 Energy ¼ QV 2
(9:2)
In the literature, the energy storage capability for a capacitor is frequently defined as follows: 1 Energy ¼ CV 2 2
(9:3)
where C is defined as the capacitance in coulombs per volt, or farads (F). This is, however, the theoretical (nominal) energy storage capability [147], which only gives an indication of the actual device performance based on the capacitance of a single electrode. For instance, in devices prepared by Arbizzani et al. [147], the actual energy delivered was typically 40% to 50% of the nominal energy storage capability. In addition, since the voltage decays when a capacitor is discharged, it is more realistic to use an average voltage ([V(charged) þ V(discharged)]=2). This assumes linear voltage decay at constant current, which is somewhat of an oversimplification. Using similar reasoning, the average power delivered by a supercapacitor is defined as: Power ¼ iV
(9:4)
where i is the average current and V is the average voltage. Discharge is usually performed at constant current, thus determining the average power requires only the monitoring of voltage during discharge.
9.3.4 Performance Implications of Supercapacitor Device Types 9.3.4.1 Type I
Current
In a Type I supercapacitor, Eox(anode) ¼ Ered(cathode), meaning that Type I supercapacitor voltages are limited by the overoxidation of the polymer, which usually occurs around 0.5–0.75 V [4]. Type I supercapacitors have a charged state in which one polymer layer is fully oxidized and the other layer is completely neutral. In the discharged state, both layers are 50% oxidized—hence at most only 50% of the total polymer’s pCathode neutralization doping capacity is used [7,143,144]. Figure 9.5 represents the charging cycle of a Type I supercapacitor: as the anode oxidizes (going from half-oxidized to fully oxidized), the cathode neutralizes (going from half-oxidized to neutral). In an idealized case, the voltage decays linearly at constant current. There is complete overlap or symmetry between the cathode and the anode in terms of oxidaAnode oxidation tion and reduction behavior, so the polarity of the device is not important. Potential Furthermore, the current will be constant at slow (50 mV=s) scan speeds at all voltFIGURE 9.5 Charging of a Type I EAP supercapacitor. ages within the operating window, hence (Adapted from Rudge, A. et al., J. Power Sources, 47, 89, 1994. a typical cyclic voltammogram of a Type I With permission. Copyright 1994, from Elsevier.)
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0.5 0.4 0.3
Current (mA)
0.2 0.1 0.0 −0.1 −0.2 −0.3 −0.4 −0.5 0.00
0.10
0.20
0.30
0.40
0.50
Potential (V)
FIGURE 9.6 Cyclic voltammogram of a Type I=II supercapacitor prepared from PEDOT=PProDOT in gel electrolyte (70% tetraethylene glycol dimethyl ether, 20% ultrahigh molecular weight PMMA, and 10% EMI-BTI) at 500 mV=s. (Reproduced from Stenger-Smith, J.D. et al., J. Electrochem. Soc., 149, A973, 2002. With permission. Copyright 2002, The Electrochemical Society.)
supercapacitor would resemble the one shown in Figure 9.6 [37]. Examples of such behavior are found in Refs. [7,37,143,144,148]. 9.3.4.2 Type II In a Type II supercapacitor, Eox(anode)– Ered(cathode) is around 0.5 V, and the overvoltage is around the same as that for a Type I, meaning that Type II supercapacitor voltages are around 1.0–1.25 V. In a charged Type II supercapacitor, the polymer with the highest oxidation potential (anode) is completely oxidized, and the polymer with the lower oxidation potential (cathode) is completely neutral. In the discharged state, the polymer with the higher oxidation potential is less than 50% oxidized, and the polymer with the lower oxidation potential is more than 50% oxidized. Thus, approximately 75% of the total polymer p-doping charge capacity is used. Figure 9.7 represents the charging cycle of a Type II supercapacitor: as the anode oxidizes (going from slightly to fully oxidized), the cathode neutralizes (going from mostly oxidized to neutral). Assuming some overlap in the current–voltage characteristics of the polymer layers and a difference in oxidation potential of 0.5 V, a typical current–voltage plot would resemble the one shown in Figure 9.6 but with a higher upper voltage limit. The general benefits of this type of supercapacitor are the increased average voltage at which the charge is delivered and the increased charge output. A typical voltage decay plot for a Type II supercapacitor is similar to that of a Type I supercapacitor but with slightly higher current and voltage. For Type II devices, the polymer with the higher oxidation potential must be used as the anode. The charge capacity of the device will be much lower in reverse bias, and stability could be a problem, because the polymer with the lower oxidation potential would be severely overoxidized [7,143,144]. 9.3.4.3 Type III and Type IV In Type III and Type IV supercapacitors, Eox(anode)–Ered(cathode) typically ranges from 0.8 to 3 V [147]. The overvoltage is approximately the same as that of a Type I supercapacitor (perhaps less due to the sensitivity of n-doped polymers), meaning that Type III and Type IV supercapacitors typically have
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Current
voltages ranging from 1.3 to 3.5 V. At this point, it is necessary to make a trade off between stability at lower voltages and Cathode neutralization higher power and energy densities. In charged Type III and Type IV supercapacitors, the polymer anode is completely oxidized, and the polymer cathode is completely reduced. In the discharged state, both polymers are neutral; hence 100% of the anode’s p-doping capacity and 100% of the cathode’s n-doping capacity are used. Anode oxidation Figure 9.8 represents the charging cycle of Type III or Type IV supercapacitor: as the Potential anode oxidizes (going from neutral to fully oxidized), the cathode reduces (going from FIGURE 9.7 Charging of a Type II EAP supercapacitor. (Adapted from Rudge, A. et al., J. Power Sources, 47, 89, 1994. neutral to fully reduced). With permission. Copyright 1994, from Elsevier.) Voltage decay of Type III and Type IV supercapacitors under constant current conditions varies markedly from that of a Type I or Type III supercapacitor (Figure 9.9). Ideally, the voltage remains high (over 2 V) until almost all the charge is passed. Figure 9.10 shows the performance of a typical Type III supercapacitor [11]. The device, prepared using poly(dithieno[3,4-b:30 ,40 -d]thiophene) as both the cathode and the anode, exhibits a capacity of 189 mC=cm2 and a coulombic efficiency of 91%. Note the difference between that cyclic voltammogram and that of a Type I supercapacitor seen in Figure 9.6.
9.3.5 Hypothetical Supercapacitor Outputs As a demonstration of the obtainable charge storage of these device types, the following hypothetical calculations are introduced. Energy and average power are calculated per kilogram of device (electrodes, electrolytes, connectors, and packaging). A 1 kg device is presumed to have 100 g of active polymer with the remaining 900 g comprising electrodes, electrolyte, separator paper, and electrical connections. Other estimations include polymer charge capacities of 150 C=g (an average and modest number for EAPs) [143], an average polymer thickness of 0.5 mm, a polymer density of 1 g=cm3 (giving specific area of 2 104 cm2=g), and the device voltages discussed earlier, the energy and power for large- and smallscale devices are predicted. The outputs predicted below are meant as general guidelines Cathode reduction and comparisons.
Current
9.3.5.1 Type I
Anode oxidation Potential
FIGURE 9.8 Charging of a Type II EAP supercapacitor. (Adapted from Rudge, A. et al., J. Power Sources, 47, 89, 1994. With permission. Copyright 1994, from Elsevier.)
Theoretically, a 1 kg device containing 100 g polymer would have a capacity of 7500 C, since only half the total capacity is used in Type I devices. Using a modest average current density of 0.5 mA=cm2 [144], a total area of 2 106 cm2, and an average voltage of 0.25 V, the average current would be 500 A, yielding an average power density of 125 W=kg of device. The energy stored in this device would be 1.9 kJ=kg. Scaling down to a small device with 1 cm2 area, the capacity would be around 3.75 mC, the energy stored would be 0.9 mJ, and the average power would be 125 mW.
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Voltage
Conjugated Polymers: Processing and Applications
Type I Type II Type III and IV
0
1 Fraction of charge
FIGURE 9.9
Comparison of voltage decay characteristics for Type I, Type II, and Type III and IV supercapacitors.
9.3.5.2 Type II Theoretically, a 1 kg device containing 100 g polymer would have a capacity of 11,250 C, assuming 75% of the total capacity is used in Type II devices. Using the same thickness as above and an average voltage of 0.375 V, the current density increases to 0.75 mA=cm2. Therefore the average current would be 750 A, yielding an average power density of 280 W=kg of device. The energy stored in this device would be 4.2 kJ=kg. Scaling down to a small device with 1 cm2 area, the capacity would be around 5.6 mC, the energy stored would be 2.1 mJ, and the average power would be 280 mW. 9.3.5.3 Type III and Type IV Theoretically, a 1 kg device containing 100 g polymer would have a capacity of 15,000 C, assuming 100% of the total capacity is used. Using the same thickness as above and the voltage ranges discussed earlier, the current increases to 1.0 mA=cm2. Therefore the average current would be 1000 A, yielding an average power density range of 1300–3500 W=kg. Energy storage would range from 19 to 52 kJ=kg. Scaling down to a small device with 1 cm2 area, the capacity would be ca. 7.5 mC. Using an average voltage of 2.25 V, the energy stored would be 17 mJ, and the average power would be 2250 mW.
10.0
I (mA/cm2)
5.0
0.0
−5.0
9.3.6 −10.0 0.0
1.0
2.0
3.0
∆V (V)
FIGURE 9.10 Cyclic voltammogram of a Type III supercapacitor prepared from PDTT in 0.2 M TEABF4 in propylene carbonate at 50 mV=s. The dotted line is the electrical response of bare carbon paper electrodes. (Adapted from Arbizzani, C. et al., Electrochim. Acta, 40, 1871, 1995. With permission Copyright 1995, from Elsevier.)
Review of Supercapacitor Literature
While their approaches may be different, the goals of most supercapacitor researchers are the same. They seek high specific capacitance to maximize energy storage, low electrical resistance to maximize power density, excellent stability, and of course low cost. The move to EAP-based supercapacitors results from several advantages of the polymer devices
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relative to the metal oxides. Charge can be stored throughout the volume of the polymeric materials, rather than just at or near the surface of the metal oxides, resulting in increased charge storage capacities. Oxidation and reduction processes in the polymer can also be very fast, if polymer morphology and electrolyte are optimized. Also, most EAPs are considerably less expensive than metal oxides and comparable in cost to most carbonaceous materials. Even with all of these advantages, EAP-based supercapacitors are not yet competitive with metal oxide supercapacitors; the best metal oxide supercapacitor, based on hydrous ruthenium oxide, exhibits a specific capacitance of 720 F=g, well above that of any polymeric system reported to date. Several reviews of metal oxide supercapacitors can be found in the literature [1–3,6,149]. The main types of polymers investigated for use in supercapacitors include polypyrrole, polyaniline, polythiophene, and polythiophene derivatives, with a few less common polymers investigated primarily for n-doping (Figure 9.4). The maximum energy and power densities of n- and p-dopable polythiophene derivatives have specific capacities (10–15 Ah=kg) that are 5–6 times higher than those of high surface area carbonaceous electrodes [150]. However, compared with cathode materials of lithium secondary batteries, such as LiCoO2 or LiNiO2 (around 100 Ah=kg), the specific capacities of current polythiophene derivatives are approximately an order of magnitude lower. 9.3.6.1 Type I Supercapacitors The bulk of EAP-based supercapacitor work to date has focused on Type I devices. Polypyrrole (PPy, Figure 9.4C) has been studied [147,151–153] for this application, with specific capacitance values ranging from 40 to 200 F=g*. Garcia-Belmonte and Bisquert [151] electrochemically deposited PPy devices that exhibit specific capacitances of 100–200 F=cm3 with no apparent dependence on film thickness or porosity; extensive modeling of impedance characteristics was used. Hashmi et al. [153] prepared PPybased devices using proton and lithium-ion conducting polymer electrolytes. As is often observed, electrochemical performance suffered somewhat in polymeric electrolytes; single electrode specific capacitances of 40–84 F=g were observed with stability of 1000 cycles over a 1 V window. Devices based on polyaniline (PANI, Figure 9.4B) have also received considerable attention [154– 159]. Prasad and Munichandraiah [156] prepared a high-energy, high-power PANI supercapacitor, exhibiting a specific capacitance of 250 F=g in a gel electrolyte. While most PANI devices are prepared from electrochemically polymerized aniline, Ryu et al. [157,158] prepared devices using chemically polymerized aniline and compared HCl, Et4NBF4, and LiPF6 dopants in acetonitrile; stability and capacitance were highly electrolyte dependent, with a maximum capacitance of 107 F=g. Addition of a solid polymer electrolyte [159] showed modest improvement. Iridium-doped PANI electrodes were found inferior to those prepared with more typical dopants [155]. Belanger et al. [154] reported just 5% loss in electroactivity of PANI electrodes over 20,000 cycles whereas the device exhibited a 33% decrease in discharge capacity. A derivative of PANI, a polydiaminoanthraquinone (PDAAQ, Figure 9.4I) oligomer, has been prepared by Suematsu and Naoi [150] to increase specific capacity relative to PANI while hopefully maintaining electrochemical stability. While PDAAQ has a theoretical specific capacity of 338 Ah=kg, the oligomer film exhibited a specific capacity of 50 Ah=kg in 4 M H2SO4; no device data were presented. Several other polymers have also been evaluated for use in Type I supercapacitors. While polythiophene exhibits an excellent specific capacitance for energy storage (250 F=g) [160], polythiophene supercapacitors are uncommon, possibly due to cycle life issues. However, derivatives of polythiophene have been used in the preparation of Type II, Type III, and Type IV supercapacitors (see Section 9.3.6.2 and Section 9.3.6.3). One class of polythiophene derivative has been studied for use in Type I supercapacitors: research into poly(3,4-ethylenedioxythiophene) (PEDOT, Figure 9.4J)-based supercapacitors has been driven by PEDOT’s superior chemical and electrochemical stability [148] as well as its fast switching times [161]. Carlberg and Inganas [148] demonstrated energy density of 1–4 Wh=kg at power densities of *Unless otherwise noted, all weights are of the polymer, not of the device.
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35–2500 W=kg. Improving energy density to at least 15 Wh=kg would meet the U.S. Department of Energy’s (DOE) electric vehicle power requirements; the authors suggest that this improvement could be met by increasing the cell voltage from 0.8 V, which they propose could be accomplished by moving to a Type II or Type III=IV device. Ghosh and Inganas [162] reported that PEDOT-based hydrogels improve ion mobility in the polymer matrix, improving power density. Alone, the hydrogel exhibits poor energy density and poor mechanical stability, but these can be improved by electrochemical deposition of PPy on the surface of the hydrogel. This process yields freestanding films with higher capacitance than standard PEDOT supercapacitors as well as extremely high power densities, even at relatively high energy densities. Sonmez et al. [161] reported the electrochemical studies of a PEDOT derivative, poly{1,3-bis[20 -(30 ,40 -ethylenedioxy)thienyl]-benzo[c]thiophene-N-200 -ethylhexyl-4,5-dicarboximide} (Figure 9.4K). This polymer can be both p-doped and n-doped, allowing for use in Type I–Type IV devices (see discussion of Type III devices in Section 9.3.6.3). They observed large, stable capacitive current in the p-doping form; the polymer can store 550 C=g in the p-doped state, with the added benefit of a 2.15 V potential window for p-doping. 9.3.6.2 Type II Supercapacitors Relatively little work has been published on Type II supercapacitors. Arbizzani et al. [147] have prepared PPy=poly(3-methylthiophene) devices; performance was similar to their PPy-based Type I device and to carbon supercapacitors. Clemente et al. [163] prepared PPy=PANI devices with specific capacitance values as high as 25 F=g, depending on electrolyte composition. Stenger-Smith et al. [37] prepared poly(3,4-ethylenedioxythiophene) (PEDOT, Figure 9.4J)=poly(3,4-propylenedioxythiophene) (PProDOT, Figure 9.4L) Type II supercapacitors. Switching speed and cycle life were found to depend heavily on electrolyte composition, with only 2% loss in capacity over 50,000 full cycles when 1-ethyl-3- methyl-1Himidazolium bis(trifluoromethanesulfonyl)imide (EMI-BTI) was used as the electrolyte. 9.3.6.3 Type III and Type IV Supercapacitors Interestingly, most of the n-doping polymers investigated for use in Type III and Type IV supercapacitors are derived from polythiophene. As part of a systematic probe of supercapacitor types, Rudge et al. [143] discussed the possibility of using other common EAPs in Type III supercapacitors. Polyacetylene (PA, Figure 9.4A) and poly-p-phenylene (PPP, Figure 9.4H) are inappropriate due to the irreproducibility of their n-doping processes whereas nonderivatized polythiophene (PT) n-doping is only reversible for very thin films [164]. Poly-3-phenylthiophenes have been extensively studied. Rudge et al. [143] compared PT, poly-3phenylthiophene (PPT, Figure 9.4F, where X ¼ H), and poly-3-(4-fluorophenyl)thiophene (PFPT, Figure 9.4F, where X ¼ F) n-doping processes. While PPT achieves higher charge density than PT, specific capacity of PPT drops off with film thickness during n-doping. PFPT, on the other hand, reaches high charge densities and undergoes reversible n-doping. The redox properties of PFPT were found to be highly dependent on electrolyte composition, with best results found for Me4NCF3SO3: Energy and power densities of 39 Wh=kg and 3500 W=kg, respectively, and cell voltages exceeded 3 V. Based on the elecrochemical data and modeling experiments, the researchers [135] proposed that the success of this material is due to electron transfer from the negatively charged polythiophene backbone back to the fluorophenyl substituent. Poly(3-methylthiophene) (PMT, Figure 9.4E) and polydithieno[3,4-b:30 ,40 -d]thiophene (PDTT, Figure 9.4M) have been prepared by Arbizzani et al. [11,144,147,165–168] and compared to PFPT for use in Type III supercapacitors. Specific capacitance values for n-doping of PFPT, PDTT, and PMT reached 80, 70, and 165 F=g, respectively, and energy densities of 7, 10, and 20 Wh=kg, respectively, with some dependence on film thickness. Based on these values as well as excellent stability over 6000 cycles and low cost, Mastragostino et al. [165] recommend PMT for Type III supercapacitors. Specific capacitance of PMT electrodes reached 240 F=g for p-doping and 180 F=g for n-doping [167]. Hybrid devices with PMT and carbon electrodes (see Section 9.3.6.5) [166,167] are also recommended.
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To obtain high energy and power densities, Ferraris et al. [169] attempted to enhance ion diffusion, since that is often the rate-limiting step in the redox process. Open rigid structures were targeted, so they prepared a series of poly-3-phenylthiophenes with fluoro, difluoro, and cyano substituents on the phenyl ring (Figure 9.4F). Film morphology was found to be dependent on the size of electrolyte anions, with larger anions increasing film porosity. Me4NCF3SO3 electrolyte solution produced more porous films with greater charge densities and faster discharge rates than films grown in Et4NBF4. While the fluoro-substituted polymer PFPT has been extensively studied by Ferraris and others, Ferraris reported that the difluoro-cyano-substituted polymers are more stable than PFPT and may yield better Type III supercapacitors. Another polythiophene derivative was synthesized by Sonmez et al. [161] for use in Type III supercapacitors. This group’s approach is based on the concept that low band gaps will increase intrinsic charge carrier densities, producing more stable n-doping polymers. Thus, they prepared poly(1,3-bis(20 [30 ,40 - ethylenedioxy]thienyl)-benzo[c]thiophene- N-200 (-ethylhexyl-4,5- dicarboximide) (Figure 9.4K). As a Type III supercapacitor, this polymer can store ca. 325 c=g of polymer over 2.15 V; the majority of the charge storage occurs via non-Faradaic processes. Limited stability studies showed minimal loss in current response over 10 cycles in Bu4NPF6=propylene carbonate; similar experiments in LiClO4=propylene carbonate were less successful, possibly due to interaction of lithium with the polymer backbone, destroying electroactivity. Loveday et al. [170] proposed the use of self-n-doping polymers in Type III supercapacitors. Self-ndoping was expected to alleviate cation transport problems and enhance the n-doping process, thus enhancing device performance. While the polymer, poly-3-(p-trimethylammoniumphenyl)bithiophene (Figure 9.4N) showed promising n-doping, no supercapacitor device properties were reported. Fusalba [171] investigated polycyclopenta[2,1-b;3,4-b0 ]dithiophen-4-one (Figure 9.4O) for use in supercapacitors. Specific capacitance values of ca. 70 F=g were demonstrated in both the p- and n-doped states, and high energy and power densities (6 Wh=kg and 1000 W=kg, respectively, for an 18 s discharge time) were obtained; cycle life was problematic. Cyclopenta[2,1-b; 3, 4-b0 ]dithiophen-4-one has been further derivatized by Ebron et al. [146] to yield 4,40 -bicyclopenta[2,1-b;3,4- b0 ]dithiophenylidene dimer. The resultant polymer (Figure 9.4P) appears to be among the most promising n-doping polymers to date, with a wide potential window (the authors expect ca. 3 V) and good cycle stability (6% decrease in charge after 1000 n-doping cycles). Fusalba [172] also prepared a series of poly(cyano-substituted diheteroareneethylenes) (Figure 9.4Q) with a wide potential range (ca. 2 V), along with energy and power densities of 8.6 Wh=kg and 1600 W=kg, respectively (for a discharge time of 20 s). Unfortunately, like many n-doping polymer devices, cycle stability was poor (60% loss over 2000 cycles). 9.3.6.4 Composite Systems In attempts to improve capacitance over that achieved to date in polymer-based supercapacitors, several groups have investigated EAP-based composites based on PPy [170,173–175], PANI [78,176–178], PT [160,179], PMT [166], and PFPT [160,180,181]. By using a combination of an EAP electrode and an activated carbon electrode, it is possible to produce devices with higher specific power than double-layer capacitors due to the lower equivalent series resistance [166]. In some cases, composite electrodes are a matter of necessity: when chemical polymerization is used to produce insoluble polymer powders, the powders can be blended with carbonaceous material and binder. Laforgue et al. [160] pointed out that electrochemically deposited polymer film thicknesses are limited. For greater storage capacities, they utilized chemical polymerization and mixed the polymers with graphite and binders to yield conductive powders that could be pasted onto current collectors. A PT=PFPT Type IV composite device was prepared in this fashion [160,180]; specific capacitances reached 260 F=g for PT and 110 F=g for PFPT, but cycle life was poor. Modification of the PFPT structure to include additional fluorine substituents did not improve performance [166]. Kim and Chung [181] prepared PFPT devices, reporting 1.6 Wh=kg and 268 W=kg energy and power densities, respectively, after 1000 cycles over a 2.7 V window. Based on interfacial resistance measurements, there
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appeared to be no loss in stability over 10,000 cycles. Hu and Chu [176] modeled equivalent circuits for PANI=graphite electrodes; they found that capacitance varied with dopant choice, with H2SO4 providing the best results. Several research groups have focused on EAP-carbon nanotube (CNT) composite electrodes, possibly because of the increased conductivity of CNTs relative to other carbonaceous materials. For a review on CNT capacitors that do not utilize EAPs, see Ref. [182]. While the EAP component was found to store charge via a pseudocapacitive redox process, the CNTs were found to store charge in the electrical double layer in traditional double-layer capacitor fashion [183]. Incorporation of the CNTs gave slight improvement in capacitance relative to analogous acetylene black-based composites [184] presumably due to increased surface area and therefore better accessibility of the polymer to the electrolyte. In the composite devices, research supports a redox-based pseudocapacitive charge storage mechanism [183], which may be enhanced by the double-layer capacitance of the CNTs. The predominant polymer used in these studies was polypyrrole [185–192], but polyaniline [190,193], poly(3-methylthiophene) [192], and poly(3,4-ethylenedioxythiophene) [184] have also been evaluated. Lota et al. [184] presented a systematic evaluation of the composite system, evaluating PEDOT and CNT electrodes separately, PEDOT– CNT composite electrodes, and PEDOT–acetylene black composite electrodes. They noted that either CNTs or acetylene black improved performance by enhancing the surface area, making the composite bulk more accessible to ions and improving electronic conduction. Researchers have shown that the composites yield higher specific capacitances than either of the individual components when used alone [187,188,193]. 9.3.6.5 Hybrid Devices Rather than preparing supercapacitors entirely from EAPs, several groups [168,179,194] have combined a p-doping EAP electrode and a carbonaceous electrode in what are referred to as hybrid devices, thus eliminating the need for a stable n-doping polymer. Polyaniline=activated carbon hybrid devices have been prepared [194]. Specific capacitance of 380 F=g was obtained, with specific energy and specific power of 18 Wh=kg and 1250 W=kg, respectively, stability over 4000 cycles, and a 1–1.6 V window. Mastragostino et al. [168] prepared hybrid devices from PMT=activated carbon. They claimed that the resultant hybrid device outperforms the double-layer carbon supercapacitors presently on the market in terms of specific energy and power. Laforgue et al. [179] prepared hybrid devices using PFPT and a variety of activated carbons. The best devices operated over a 3 V window, with energy density and power density of 48 and 9000 W=kg, respectively, and specific capacitance up to 68 F=g with stability over 8000 cycles. 9.3.6.6 Electrolyte Considerations As was noted repeatedly above, electrolyte choice can have a significant effect on device performance. For instance, use of lithium cations in electrolytes for n-doping polymers may result in insertion of lithium into the polymer backbone, destroying electroactivity [161]. Electrolyte choice during electropolymerization can also significantly affect polymer morphology, which in turn can affect charge density and discharge rates [169]. To function effectively in a supercapacitor, electrolytes must tolerate the voltage ranges, be unreactive with the polymers, and have good ionic conductivity [8,195]. Salts frequently used in supercapacitor electrolytes include LiClO4, LiCF3SO3, Li- BTI, and R4NBF4. The wide variety of EAPs, electrolytes, and solvents leads to the conclusion that the ion-transfer processes for EAPs in any electrochemical reaction are complex interactions among ions, the polymer film, and the solvent (if any). Any type of ion transfer is permissible as long as conservation of charge is obeyed. There is evidence in the literature [14,148,196–198] that both the cation and the anion move into and out of the active polymer layer during redox processes. However, using a polymeric anion with a small molecule cation virtually guarantees that only the cation will move in and out of the polymer layer and vice versa [199]. Ion migration is also accompanied by solvent migration. This factor is important in choosing an electrolyte system, since the solvent shell of an ion (such as lithium) can be large compared to that of the ion.
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Most supercapacitors have been prepared using electrolytes in solution. For polymer stability reasons, early aqueous work gave way to organic electrolytes, frequently acetonitrile, ethylene carbonate, or propylene carbonate. While these solutions are usually highly conductive, they are subjected to leakage and evaporation problems. While moving to higher boiling solvents reduces evaporation, it also increases viscosity and therefore reduces conductivity. All-solid-state supercapacitors are desirable to prevent or reduce corrosion, self-discharge, low energy density, and bulky design [153]. The two types of electrolytes used in all-solid-state supercapacitors are gel electrolytes and solid polymer electrolytes (typically used in battery and fuel cell membranes). Ion conducting solid polymer electrolytes, such as those used in battery and fuel cell membranes, have been explored for use in supercapacitors [153,159,200,201]. While these electrolytes are environmentally benign and do not leak, conductivities are typically much lower than liquid or gel electrolyte systems, especially at subambient temperatures (important for military and space applications). Nevertheless, capacitance in supercapacitors prepared with solid polymer electrolytes has been reported to be as good as or better than the same devices constructed using liquid electrolytes. Nafion [200], polyethylene oxide [153], and polyvinyl alcohol [153] are the polymers of choice for this application. Gel electrolytes are typically comprised of electrolyte, solvent, and a polymer, usually poly(methyl methacrylate) [163,201–204] or poly(ethylene glycol) [205]. The increased viscosity of these electrolytes yields conductivities intermediate to those of solution electrolytes and ion conducting solid polymer electrolytes. Chojnacka et al. [206] studied electrolytes combining LiCF3SO3 and a polymer in ethylenecarbonate=gamma-butyrolactone. The resulting gel electrolytes exhibit high ionic conductivity and good electrochemical stability. Prasad and Munichandraiah [156] used a gel electrolyte in a PANI-based Type I supercapacitor and achieved specific capacitance and specific power greater than anything reported in the literature for traditional electrolyte systems. The low volatility, nonflammability, and high ionic conductivity of ionic liquids (room temperature molten salts) have led to recent studies to improve the high-temperature safety and durability of a variety of EAP-based electrochemical devices including supercapacitors. Long-term device stability is typically enhanced when replacing traditional electrolyte solutions with ionic liquid electrolytes because of the very low volatility of the ionic liquids [34–36,207,208]. An electrochromic device was cycled 1 million times with negligible loss in performance [207]. Replacing traditional organic solvent-based electrolytes with ionic liquid electrolytes is also hoped to improve energy density and thermal stability of supercapacitors [209,210]. Sato et al. [209] reported on a series of ionic liquids for use as capacitor electrolytes. In comparison with traditional organic liquid electrolytes such as tetraethylammonium tetrafluoroborate in propylene carbonate, carbonaceous double-layer capacitors utilizing an ionic liquid electrolyte, (N,N-diethyl-N-methyl-N-(2-methoxyethyl)ammonium tetrafluoroborate), exhibit a higher capacity above room temperature and better cycle life at 1008C (but performance is not as good below 408C). Stenger-Smith et al. [37] prepared Type II supercapacitors using gel electrolytes based on both lithium bis(trifluoromethanesulfonyl)imide (Li-BTI) and an ionic liquid, 1-ethyl- 3-methyl-1H-imidazolium bis(trifluoromethanesulfonyl)imide (EMI-BTI). While devices using both electrolyte systems switch rapidly and store similar amounts of charge, the ionic liquid electrolyte provided significant cycle life improvement, with the device losing only 2% of its charge capacity over 50,000 cycles relative to a 30% loss in charge capacity for the Li-BTI device under the same conditions. Supercapacitor electrolytes will be a critical factor in constructing working Type III and Type IV supercapacitors. The ion movement processes at both the n-doped and p-doped polymer electrodes play a critical role in the functioning of the device. There is ample data concerning the effects of the electrolyte on p-doping polymers; some electrolytes work quite well allowing fast switching time and tens of thousands of deep cycles whereas others give quite the opposite result; behavior is severely compromised if an EAP is grown in one type of electrolyte and cycled in a different type of electrolyte. While this may be true for n-doping polymers, the severe shortage (or complete lack) of true n-doping polymers hampers progress in this area. As n-doping polymers continue to be studied, it is possible that a whole new type of electrolyte, or a hybrid electrolyte consisting of two types of salt (for e.g., for a Type IV supercapacitor) may be necessary.
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Critical Analysis of Supercapacitor Research
In order to standardize the results from the literature, the following suggestions are made. In future publications, it should be made clear whether analysis is of a single electrode, or of a complete device. True EAP polymer supercapacitor cells contain exactly two electrodes and a thin layer of electrolyte. It is also important to report the area of the cell, the spacing between the active layers, and if possible, to estimate the energy and power density based upon the weight of the entire device, not just the active layer. The speed of future devices should also be clearly stated. Reports of the number of cycles should include the percent DOD. Many electrochemical devices will last for 106 cycles at 1% DOD. In order for a device to be practical, and for the testing to reflect the real application, DOD during cycling should be at least 60%. There is a great supply of p-doping polymers (PPy, PT and derivatives, and polyalkylene dioxythiophenes) with adequate charge capacity and stability. The development of a new p-doping material is not critical to the development of an EAP-based supercapacitor. The most pressing need is for true n-doping polymers to be synthesized. Most voltammetric analyses of reduction processes in conducting polymers show very sharp well-defined processes and none of the pseudocapacitive behavior seen in most (if not all) p-doping polymers. This begs the question as to whether the process is truly n-doping rather than just charge trapping. On top of this requirement, the n-doping process in the polymer must be extremely stable not only during cyclic voltammetry but also in a device. For practical devices containing n-doping polymers, some type of hermetic sealing will likely be necessary to prevent exposure to oxygen and water.
9.4
Conclusions
In summary, a wide range of EAPs has been used to demonstrate the capability and capacity of polymerbased charge storage devices. While the differences between EAP-based batteries and supercapacitors may be small, it is clear that certain applications require either a battery or a supercapacitor and that such devices can be based upon EAPs. This review of the literature reveals some promising results for both EAP batteries and supercapacitors. Many EAPs excel as anode materials (oxidation neutralization), and there are a few EAPs that might be used as cathode materials (reduction neutralization). At present, it may be possible to construct commercially viable EAP batteries and supercapacitors, but at the current state of the art it will be necessary to use some other type of nonpolymeric material for the cathode (although these materials also have drawbacks). It is therefore crucial that stable, reversible n-dopable polymers be synthesized in quantity and studied. Furthermore, there are many promising electrolytes being researched, and it is likely that electrolyte formulations will be different for each application. While traditional primary batteries and capacitors will not be obsolete any time soon, the increased demand for portable power will create new markets for rechargeable energy storage devices. It is anticipated that the size, shape, weight, capacity, energy density, and power density of these new devices will exceed current capabilities.
Acknowledgments The authors wish to thank Dr. Michele Anderson and Dr. Paul Armistead at the Office of Naval Research for continued support.
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10 Conjugated Polymer–Based Photovoltaic Devices 10.1 10.2
Introduction..................................................................... 10-1 Device Architectures and Nanomorphology Optimization.................................................................... 10-5 Bilayer Heterojunction Devices Devices
10.3
.
Bulk Heterojunction
Photon Harvesting via Band Gap Engineering........... 10-11 Limitations to Power Conversion Efficiency . Optical Modeling of Thin-Film Plastic Solar Cells . Low Band Gap Polymers
10.4
Charge Transport and Recombination ........................ 10-14 Charge Carrier Mobility . Charge Mobility and Recombination Studied by the Photo-CELIV Technique Reduced Recombination in the P3HT and PCBM Bulk Heterojunction
10.5
Contacts ......................................................................... 10-24 The Metal–Insulator–Metal Structures Voc in Bulk Heterojunction Solar Cells
10.6
A.J. Mozer and N.S. Sariciftci
10.1
.
.
Origin of the
Device Models of Bulk Heterojunction Solar Cells.... 10-27 The Equivalent Circuit Model . Extended One-Diode Model . Electric Field–Dependent Dissociation of the Coulomb-Coupled E–H Pairs . Numerical Solution to the Drift-Diffusion Equations
10.7
Concluding Remarks and Outlook.............................. 10-30
Introduction
Why to use conjugated polymers? Conjugated polymers combine the low-cost, high production-yield processability of conventional plastics with the characteristic electronic features of organic semiconductors. In addition, we list four reasons why conjugated polymers are considered promising materials for fabricating cost-effective, high-performance photovoltaic devices. 1. Conjugated polymers are in principle intrinsically stable upon photoexcitation with visible light. The electronic structure of p-conjugated polymers is schematically illustrated in Figure 10.1. Their primary structure (backbone) is determined by the s-bonds formed between the carbon atoms and various heteroatoms, for e.g., hydrogen, sulfur, and nitrogen. The bonding p-band and antibonding p*-band, on the other hand, are formed by the delocalization of the formerly unpaired pz-electrons of each carbon atoms. The energy gap between the top of the p-band 10-1
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σ*-band
π*-band formed by linear combinations of pz orbitals
∆E 1.5 eV–2.5 eV π-band
σ and π bands often overlap σ-band
FIGURE 10.1
Schematic band diagram of p-conjugated polymers.
(highest occupied molecular orbital, HOMO) and the bottom of the p*-band (lowest unoccupied molecular orbital, LUMO) is typically between 1.5 and 2.5 eV. Absorption of a photon of the visible part of the electromagnetic radiation (the wavelength range that is the most interesting if solar light harvesting applications are considered) promotes an electron from the completely filled p-band to the p*-band. The energy required for exciting electrons from bonding to antibonding s-orbitals, however, is found at much higher energies of the electromagnetic spectrum typically in the UV region. Excitation of conjugated polymers with visible light therefore leaves the primary structure formed by the s-bonds intact, which should make them intrinsically stable against photodegradation in an inert atmosphere. 2. High absorption cross section for photon harvesting. Ideally, the electronic structure of conjugated polymers is that of a one-dimensional semiconductor in which the absorption coefficient increases steeply above the band gap absorption. The absorption coefficient of the frequently used poly(p-phenylene vinylene) polymer MDMO–PPV (Figure 10.2.) reaches 105 cm1 just H3CO C6H13
OR
n S
n
P3HT
R = C10H21; MDMO-PPV R = C8H17; MEH-PPV
O
OMe O
C60[PCBM]
Me
O
C70[PCBM]
FIGURE 10.2 Chemical structures of materials commonly used in bulk heterojunction solar cells. MDMO-PPV, poly[2-methoxy-5-(30 ,70 -dimethyloctyloxy)-p-phenylene vinylene]; MEH-PPV1 poly[2-methoxy-5(2-ethylhexoxy)1, 4-phenylene vinylene]; P3HT, poly(3-hexylthiophene); C60[PCBM], 30 -phenyl-30 H-cyclopropa[1,9][5,6]fullereneC60-Ih-30 -butanoic acid methyl ester; C70[PCBM], 30 -phenyl-30 H-cyclopropa[8,25][5,6]fullerene-C70-D5h(6)-30 butanoic acid methyl ester.
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0.25 eV above its band edge. In contrast, the absorption coefficient of crystalline silicon at 300 K reaches 105 cm1 at 1.88 eV above its band edge, which in turn means that only a few hundred nanometer thin films of conjugated polymers absorb light efficiently. As it will turn out to be, this high cross section for photon absorption in conjugated polymers is a crucial prerequisite for the fabrication of efficient photovoltaic devices, which typically exhibit low charge mobility, resulting in small drift and diffusion distances. 3. Tunable bandgap within the entire visible spectral range. The minimum energy required for exciting an electron from the p- to the p*-band, i.e., the optical bandgap, is a sensitive function of the conformation of the backbone, the electron donating or withdrawing effects of various substituents, the bond length alternation, aromaticity, and interchain interactions, and can be varied using a set of guidelines often referred to as bandgap engineering [1]. The absorption spectrum of conjugated polymers therefore can be tuned to match the entire solar spectrum, or well-defined portions of it. 4. High yield of charge generation when mixed with electron acceptor materials. In small molecule polycrystalline or amorphous organic solids, it is generally assumed that the charge generation mechanism follows the Onsager model [2]. Due to the low mobility of the photogenerated charges, the mean free path is shorter than the Coulomb radius of the electron–hole (e–h) pair (weak dielectric screening) [3]. Therefore, the photoexcited e–h pairs are attracted by their mutual Coulomb field, and additional energy supplied in a secondary process is required for generating free charge carriers [4]. Such a charge generation mechanism has been challenged in one-dimensional conjugated polymers. It is predicted that the strong coupling of the fundamental photoexcitations to the lattice (electron–phonon interaction) facilitates charge separation against the attractive Coulomb (electron–hole) interaction [5]. The Su–Schrieffer–Heeger model [6] can explain the various unique photophysical features of conjugated polymers, namely the unusual charge–spin correlation of the photoexcitations [7] (solitons, polarons, bipolarons, etc.), the infrared-active vibrational modes, the characteristic photoinduced absorption of neutral and charged photoexcitations, and the strongly anisotropic line shapes of the absorption and electroabsorption spectra of stretched oriented PPV samples [8]. It appears that interchain coupling strongly determines the branching ratio between prompt luminescence (PL, radiative deexcitation of the initially excited singlet exciton) and photoconductivity (PC, electron and hole separated, and collected at external electrodes). Recent estimates in substituted PPV derivatives is 90:10, meaning that 10% of the photoexcitations lead directly to charge separation within 100 fs after initial photoexcitation and the remaining photoexcitations immediately form bound electron–hole pairs (excitons) [9]. Extensive work in the early 1990s showed that the above branching ratio is unfavorable for photovoltaic applications. It can, however, be drastically altered by mixing conjugated polymers with strong electron acceptors. Conjugated polymers with high electron affinity [10], small molecule electron acceptors [11], nanostructured metal oxides such as TiO2 [12] or ZnO [13] and semiconducting nanoparticles [14] have been demonstrated as suitable electron acceptors in conjunction with conjugated polymers as donors. The composites of these donor and acceptor materials exhibit a pronounced photoinduced electron transfer that takes place between the photoexcited state of the conjugated polymer and the electron acceptor. On photoexcitation of the conjugated polymer, a cascade of reactions is initiated leading to full charge separation [15]: D þ A !1,3 D þ A 1, 3
D þ A !1,3 ðD AÞ
(excitation on donor)
(10:1)
(excitation delocalized on the donor---acceptor complex, i:e:, exciplex) (10:2)
1, 3 ðD AÞ !1,3 Ddþ Ad
(partial charge transfer)
(10:3)
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Ddþ Ad !1,3 ðDþ A Þ 1, 3
ðD þ A Þ ! Dþ þ A
(ion---radical pair formed) (charge separation)
(10:4) (10:5)
The different steps can relax back to the initial state by either radiative (PL) or by nonradiative recombinations. The final stage of full charge separation is a necessary condition for realization of photovoltaic devices. The experimental evidence of photoinduced electron transfer between conjugated polymers and C60 using various spectroscopic tools has been discussed in detail in Ref. [15], and is briefly summarized below, highlighting some of the recent experimental findings: 1. The absorption spectrum of the conjugated polymer MDMO–PPV (Figure 10.2) blended with PCBM is a simple superposition of the absorption spectrum of the two components, and no indication of ground state interaction is observed. The absorption coefficient of MDMO–PPV, PCBM, and their 1:4 mixture by weight, calculated from their respective dielectric functions, is shown in Figure 10.3. 2. Instead of the single triplet–triplet absorption band centered around 1.35 eV in the photoinduced absorption (PIA) of pristine MEH–PPV, a sharp PIA edge at 1.1 eV and a plateau at 1.6–2.0 eV is observed in the MEH–PPV:C60 blends [16]. 3. Using photoinduced absorption detected magnetic resonance (PIADMR) [17], the 1.35 eV peak in the pristine polymer has been attributed to triplet–triplet absorption, which is quenched in the photovoltaic blends with C60. The photoinduced charge transfer is sufficiently fast to quench singlet to triplet intersystem crossing in MEH—PPV, resulting in absorption features attributed to positive radical cations (photodoping). 4. Subpicosecond transient absorption studies confirmed that the photoinduced charge transfer happens within <45 fs after photoexcitation [18]. The observed ultrafast timescale of electron transfer ensures a nearly 100% quantum yield for photoinduced charge separation at the electron donor–electron acceptor interface. 5. Steady-state PC experiments showed that blending conjugated polymers with only 1% C60 increases the PC by orders of magnitude. Transient photoconductivity experiments concluded that the increased steady-state photoconductance is the cumulative effect of higher charge generation yield and the increased lifetime of the photoinduced charges [19].
Abs. coeff. a (cm−1)
2.0 105
1.5
MDMO-PPV PCBM Blend 1:4
105
1.0 105
5.0 104
0.0 300
400
500
600
700
800
900
Wavelength l (nm)
FIGURE 10.3 Absorption coefficients of MDMO-PPV (short dashed line), PCBM (dashed line), and the MDMO– PPV: PCBM 1:4 mixture by weight (solid line) calculated from their respective dielectric functions. (Reprinted from Hoppe, H., Arnold, N., Sariciftci, N.S., and Meissner, D., Sol. Ener. Mater. Sol. Cells, 80, 105, 2003. With permission. Copyright 2003. Elsevier.)
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6. Definitive evidence for the occurrence of photoinduced charge transfer was obtained by lightinduced electron spin resonance (LESR) studies [11,20]. On photoexcitation of the MEH– PPV:C60 mixture with energies larger than the p – p* gap of the conjugated polymer, two LESR signals are detected. The signal around g < 2 is characteristic of a C60 radical anion alongside with the g 2.002, which is typical of a conjugated polymer radical cation (positive polaron). Time-resolved LESR studies also showed two kinds of photoinduced radicals at liquid helium temperatures: one prompt component decaying within nanoseconds after switching off the light and another quasistable persistent radical, which remains detectable for hours indicating persistent photoconductivity at low temperatures [21]. 7. Photoluminescence and photocurrent detected magnetic resonance studies showed that at least some portion of the initially formed charges in the MDMO–PPV and PCBM mixture form coulombically coupled radical pairs at lower PCBM concentrations. It is expected that due to the low dielectric constant of organic materials, the photogenerated electron–hole pairs are bound by their mutual Coulomb attraction. As the PCBM concentration is gradually increased above 25%, however, no indication of the formation of such pairs is observed [22]. Several models have been proposed to explain the high yield of free charge generation, which includes (a) a disorder-assisted charge separation of the Coulomb-coupled pairs: The energy gain by electron transfer to adjacent, lower lying energy sites compensate for the Coulomb interaction as argued by Offermans et al. [23]; (b) the presence of oriented dark dipoles, which localizes the hole on the conjugated chain adjacent to the electron acceptor and facilitates the separation of the Coulomb-coupled pair against geminate recombination [24]. 8. Wienk et al. have introduced a novel electron acceptor C70[PCBM] that has stronger absorption in the visible range up to 700 nm as compared to C60[PCBM] [25]. The photovoltaic devices based on the MDMO–PPV: C70[PCBM] mixture showed an indeed improved external quantum efficiency as compared to the MDMO–PPV:C60[PCBM] mixture. Using photoinduced absorption and transient absorption in the subpicosecond time domain, essentially the same charge separated state is obtained with the same forward electron transfer dynamics (experimental resolution 500 fs) independent of whether the C70[PCBM] is selectively excited at 660 nm, or both MDMO–PPV and C70[PCBM] at 510 nm. This experimental finding confirms previous results that light absorption by both the conjugated polymer and the electron acceptor, the latter termed as photoinduced hole transfer, leads to charge generation [26]. In summary, conjugated polymers are studied for high-performance, cost-effective photovoltaic applications, when mixed with suitable electron acceptors. The charge generation yield at the donor– acceptor molecular interface approaches unity, and the lifetime of the photogenerated charges is relatively long. In the following, we shall discuss the critical parameters of conjugated polymer–based photovoltaic devices in detail.
10.2
Device Architectures and Nanomorphology Optimization
10.2.1 Bilayer Heterojunction Devices The organic solar cell exhibiting the benchmark power conversion efficiency of around 1% was demonstrated by Tang in 1986 [27]. That cell was based on a double-layer-structure of Me-Ptcdi (N,N 0 -dimethyl-perylene-3,4,9,10-dicarboximide)=ZnPc (zinc-phthalocyanine), both of which are rather inefficient photovoltaic materials when used in single layer structure. The observed high performance of that bilayer heterojunction is attributed to the efficient charge generation at the interface between the two organic materials with different electron affinity. Figure 10.4 illustrates the energy diagram of two large band gap, intrinsic semiconductors before making a contact. It is expected that such a bilayer heterojunction inserted between a high work-function electrode (El1) matching the highest occupied molecular orbital (HOMO) level of the donor and a low work-function electrode
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ee 0 −ej
x AD
AA
EGD El2 EGA
El1
FIGURE 10.4 Schematic band diagram of two large bandgap intrinsic semiconductors with different electron affinity before making contact. The electron affinities (AD and AA are the electron affinity of the donor and the acceptor, respectively) are defined versus the electron energy in vacuum at the same electrical potential. EGD and EGA is the band gap energy of the electron donor and electron acceptor, respectively.
(El2) matching the lowest unoccupied molecular orbital (LUMO) level of the electron acceptor should function as a diode with a rectifying current–voltage characteristics. Under the forward bias (the low work-function electrode is biased negative in respect to the high work-function electrode) the electron injection into the LUMO of the acceptor layer from the low work-function electrode as well as the electron extraction out of the HOMO of the donor by the high work-function electrode is energetically possible and a high current can flow through the bilayer when the injected charges can move in their respective thin films. Under reverse bias (the low work-function electrode is biased positive in respect to the high work-function electrode), the electron removal from the electron donor and electron injection to the electron acceptor is energetically unfavorable. These considerations have been proposed by Aviram and Ratner over 20 years ago for a single molecule with donor-bridge-acceptor type molecular structure [28]. On illumination, the excitons generated within the exciton diffusion length from the donor–acceptor interface can reach the interface within their respective exciton lifetime, where charge separation is favorable when the following condition is fulfilled: ID AA UC < 0
(10:6)
where ID is the ionization potential of the photoexcited donor, AA is the electron affinity of the acceptor, and UC denotes the total Coulomb correlation energies. The separated charges, which overcome their mutual Coulomb field, can move away from the interface and be collected selectively at the electrodes. Bilayer heterojunction devices have been fabricated based on a spin-coated layer of conjugated polymer donors and different acceptors such as evaporated C60 layers [29–32] or high electron-affinity conjugated polymers [33–35] sandwiched between ITO coated glass and evaporated gold electrodes. In the MEH-PPV and C60 bilayer devices, the current–voltage characteristics show an exponential turn-on up to 500 mV in forward bias, and the rectification ratio is 104. The current–voltage characteristic of the device changes dramatically on illumination by visible light and performs photovoltaic effect with an open circuit voltage (Voc) around 0.5 V and a short circuit current density (Jsc) is 2 mA=cm2 with 10 mW illumination, resulting in a power conversion efficiency of 0.04% [29].
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Several reports have appeared in the recent years on improved devices using bilayer heterojunctions [36]. An important step toward superior performance was achieved by tuning the maximum of the optical field to near the bilayer heterojunction [37,38]. The insertion of exciton blocking layers [39], or chemically doped charge transport layers (so called p-i-n diode structures) [40], lead to further improvements and there is a great reproducibility in device production. However, the photoactive layer is limited to the geometrical size of the donor–acceptor interface within a thickness of around 10 nm (exciton diffusion length). Therefore, the photocurrent response is strongly limited. To improve the photoresponse, the enhancement of the donor–acceptor interfacial area is needed. This has been accomplished by using the entire volume of the device (bulk) as donor–acceptor heterojunction [10,41].
10.2.2 Bulk Heterojunction Devices The efficiency of converting the photons incident on the sample to electrons extracted to the external load hipce can be conceptually divided to the individual efficiencies of light absorption habs, exciton diffusion to the narrow interface region between the electron donor and electron acceptor hexc, the efficiency of charge generation hgen, and charge extraction hextr , so that hipce ¼ habs hexc hgen hextr . The problem with the organic photovoltaic devices such as shown in Figure 10.4 is the rather short exciton diffusion length. Consequently, most of the excitations created in the bulk will relax radiatively or nonradiatively before reaching the donor–acceptor interface. To ensure that all photogenerated excitons reach a donor–acceptor interface, the heterojunction formed between the two materials has to be scaled down to the nanometer level to form an architecture that is referred to as bulk heterojunction [42]. As such, the bulk heterojunction can be regarded as an ensemble of nanoscale heterojunctions distributed all over the volume forming a bicontinous network. A definitive advantage of a bulk heterojunction is that it can be formed by simply mixing the donor– acceptor materials in a common solvent, and cast with well-known solution deposition techniques, such as spin coating [43], the doctor blade technique [44], screen printing [45], or an evaporative spray deposition technique [46]. The details of forming a bicontinous, interpenetrating network of the two materials that enable both efficient charge generation and extraction of the charges to the electrodes, however, turned out to be a challenge. Shaheen et al. have shown that by changing the solvent from which the MDMO–PPV=PCBM active layer is cast from toluene to chlorobenzene, the hipce is increased by almost 50%, and the power conversion efficiency is more than the double up to 2.5% [43]. Extensive nanomorphological studies have revealed that the improved conversion efficiency of the chlorobenzene cast device is mainly due to the finer (<20 nm) nanoscale phase separation of the electron donor–electron acceptor blend as compared to the toluene-cast ones (>80 nm grain sizes [47– 51]. One interesting question addressed by van Duren et al. is why the best performance of the chlorobenzene cast device is found around 1:4 ratio [48]? The contribution of PCBM to the total absorption is significantly smaller in the visible spectral range than that of MDMO–PPV; therefore, by increasing the PCBM content, the fraction of absorbed photons is decreasing at a constant film thickness. Moreover, charge transport is less efficient in pure MDMO–PPV as compared to PCBM [52]. It was anticipated that by diluting the two materials, the charge mobility further decreases. Van Duren et al. have proposed that in the best performing 1:4 ratio of MDMO–PPV:PCBM devices, spin cast from chlorobenzene, a hierarchical built-up of two cooperative interpenetrating networks, phase separated on two different length scales are present. The primary structure is a well-intermixed MDMO– PPV–PCBM phase corresponding to roughly 1:1 ratio of the two materials, in which virtually all photogenerated excitons reach the donor–electron interface as indicated by the efficient quenching of the MDMO–PPV luminescence. In addition, a larger scale (up to 100 nm) phase separation is observed. The larger phase is attributed to PCBM crystallites that can serve as highly mobile channels (highways) for the electrons toward the electrodes. Hoppe et al. claimed that in addition, small nanospheres attributed to coiled polymer chains are also present throughout the films [49], which may serve as the conduction channel for the positive charge. Unfortunately, the phase separation in the MDMO– PPV=PCBM blend is unstable, and increases dramatically at elevated temperatures [53]. The PCBM
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tends to form micrometer size, well-ordered crystals and diffuse out from the intermixed blend. As a result, the photoluminescence of MDMO–PPV can be detected, and a significant degradation of the device performance is observed. To overcome this morphological instability, cross-linkable donor or acceptor molecules are proposed. The results of preliminary studies show that cross linking of the vinyl substituted donor or acceptor molecules fixes the prearranged nanomorhology in the bulk heterojunction, thus preventing the morphological instability [54]. A relatively stable morphological phase separation has been achieved using a regioregular poly(3-hexylthiophene) P3HT=PCBM blend [55]. If deposited from chloroform, the films show a blue shifted and broadened absorption characteristic of amorphous-like P3HT. On postproduction treatment [56], however, the absorption spectrum is dramatically changed, featuring an overall red shift of the absorption onset and a more structured line shape characteristic of well-ordered regioregular P3HT. The hipce as shown in Figure 10.5 is increased by a factor of 2–3 at almost all absorption wavelengths. Optical modeling showed that the increased absorption accounts for roughly 40% of the total increase in the short circuit current [57]. The additional improvements and the high filling factor (FF) have been attributed to improved transport of the photogenerated charges. In overall, the power conversion efficiency of such systems is around 4%–5% [58–60]. A detailed study employing transmission electron microscopy (TEM) and selected area electron diffraction (SAED) showed that the morphology of the P3HT=PCBM films before heat treatment consists of a homogenously intermixed, rather amorphous P3HT or PCBM phase (Figure 10.6a and Figure 10.6b). In addition, fibrillar P3HT crystals with 15 nm diameter and <500 nm length are also observed [55]. On heat treatment, the fibrillar P3HT crystals dominate throughout the whole film and form an interconnected network (Figure 10.7a and Figure 10.7b). This finding is supported by previous reports on the two-dimensional stacking of regioregular P3HTmacromolecules, resulting in two-dimensional lamellar structures [61,62].
70 60
IPCE (%)
50 40 30 20 10 0 400
450
500
550
600
650
700
Wavelength (nm)
FIGURE 10.5 Incident photons to converted electrons (IPCE) measured for a P3HT=PCBM solar cell as prepared by the spin coating technique from 1,2-dichlorobenzene solutions (open triangles), after thermal annealing (open squares), and simultaneously treated by thermal annealing and an external voltage (filled circles). (Reprinted from Padinger, F., Rittberger, R.S., and Sariciftci, N.S., Adv. Func. Mater., 13, 1, 2003. With permission of Wiley–VCH, Weinheim, Germany.)
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(a)
Fibrillar-like P3HT crystals
PCBM nanocrystals in the matrix (b)
FIGURE 10.6 (a) Bright-field transmission electron microscopy (TEM) images of pristine (without heat treatment) photoactive layer of P3HT=PCBM composite film solar cells and (b) its corresponding schematic representation. The inset (in Figure 10.6a) shows the small area electron diffraction pattern. (Reprinted from Yang, X., Loos, J., Veenstra, S.C., Verhees, W.J.H., Wienk, M.M., Kroon, J.M., Michels, M.A.J., and Janssen, R.A.J., Nano Lett., 5, 579, 2005. With permission of the American Chemical Society.)
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(a)
(b)
FIGURE 10.7 (a) Bright-field transmission electron microscopy (TEM) images of the annealed (1208C for 60 min) photoactive layer of P3HT and PCBM composite film solar cells and (b) its corresponding schematic representation. The inset (in Figure 10.7a) shows the small area electron diffraction pattern. (Reprinted from Yang, X., Loos, J., Veenstra, S.C., Verhees, W.J.H., Wienk, M.M., Kroon, J.M., Michels, M.A.J., and Janssen, R.A.J., Nano Lett., 5, 579, 2005. With permission of the American Chemical Society.)
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The P3HT=PCBM bulk heterojunction device remained stable for 1000 h with less than 20% change in the device parameters tested at elevated temperatures, which indicates that both the chemical stability of the materials and the stability of the phase-separated network are promising for long-term operation [63]. Using encapsulation techniques, the stability of the devices can be improved to beyond 3000 h [64].
10.3
Photon Harvesting via Band Gap Engineering
10.3.1 Limitations to Power Conversion Efficiency The efficiency of converting solar to electrical energy by a solar cell depends on the band gap of the lightabsorbing semiconductor [65]. The fraction of absorbed photons of the solar photon flux increases with decreasing band gap, and reaches maximum when the band gap is zero. The open circuit voltage, which is determined by the maximum splitting of electrochemical energy of the electron–hole pair, is zero when the band gap is zero, and increases with the increasing band gap. The power conversion efficiency, therefore, shows a maximum as a function of the band gap. Calculations in 1961 by W. Shockley and H.J. Queisser showed that the power conversion efficiency of a pn-junction solar cell reaches its maximum at 30% when the band gap is around 1.1 eV [66]1. The main loss mechanisms leading to this maximum theoretical power conversion efficiency are attributed to 1. unabsorbed photons of the polychromatic light at energies lower than the band gap of the semiconductor, 2. loss of energy of the electron–hole pair due to thermalization, 3. thermodynamic losses in converting the chemical energy of the electron–hole pair into entropyfree electrical energy, and 4. device electrical losses that are usually defined by the filling factor (FF) [67]. Most of the above loss mechanisms play a role in conjugated polymer bulk heterojunction solar cells also, as illustrated in Figure 10.8. First, due to the rather large band gap of commonly used conjugated polymers, a significant portion of the incoming photons are not absorbed at all (process 1). Second, absorption of a photon with energies higher than the band gap excites the molecule to higher lying electronic and vibronic excited states. The excitation rapidly relaxes to the lowest electronic excited state by emitting phonons (thermalization [process 2]). If an electron acceptor is present, the ultrafast photoinduced charge transfer may compete with the thermalization of the excited state leading to charge separation (process 3 in Figure 10.8) [18].
E
Donor
Acceptor LUMO+1
2 LUMO (~ ID*)
hn < Eg 1
Eg
ket
3
∆G°et < 0
LUMO(~AA ) eVoc
HOMO HOMO
FIGURE 10.8 The energy diagram of the photoinduced electron transfer and the main energy loss mechanisms in donor–acceptor bulk heterojunction solar cells.
1
Calculated assuming only radiative recombination and that the absorptivity approaches 1 for wavelengths shorter than the band gap absorption.
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The driving force for the electron transfer is the gain in free energy DGet0 < 0. From Figure 10.8, it is evident that charge generation by photoinduced electron transfer lowers the energy of the electron, which energy is dissipated by emitting phonons. As a consequence, the photoinduced electron transfer lowers the splitting of the electrochemical energy of the electron and hole pair, thus decreasing the maximum theoretical limit of the open circuit voltage. The routinely observed 0.8 V open circuit voltage of bulk heterojunction solar cells using MDMO–PPV=PCBM is generated by the absorption of a photon with the energy of at least 2 eV. On the basis of Figure 10.8, maximizing the power conversion efficiency of a bulk heterojunction solar cell involves the optimization of all the four relevant electronic levels, i.e., the HOMO and LUMO of the conjugated polymer donors and the HOMO and LUMO of the acceptor. First, a low band gap conjugated polymer is preferred so that its absorption better matches the solar spectrum, as will be shown later in this section. Second, the free energy gain of electron transfer DGet0 should be minimized by decreasing the energy offset between the conjugated polymer LUMO and the acceptor LUMO. It is important that the fast photoinduced electron transfer rate ket , prerequisite for high quantum yield of charge separation at the donor–acceptor interface, depends on DGet0. According to the Marcus theory of electron transfer, the electron transfer rate is proportional to the sum of free energy gain and the reorganization energy l as [68] ket ¼ K exp (DGeto þ l)2 =4lRT ,
(10:7)
where K is constant that includes a probability factor and the electronic coupling between the donor and the acceptor [69]. Ideally, the electron transfer rate should be optimized so that it is still faster than any of the exciton decay channels (which must include exciton transport to the donor–acceptor interface), yet the free energy gain upon electron transfer is lowered. Lastly, as discussed in more detail in Section 10.5, the position of the HOMO level of the conjugated polymer determines the maximum open circuit voltage on the hole contact side, and should be lowered.
10.3.2 Optical Modeling of Thin-Film Plastic Solar Cells The typical layer-stack of a bulk heterojunction solar cell is shown in Figure 10.9a. The total thickness of the photoactive layer of most devices is less or comparable to the penetration depth of the visible light (<500 nm). Therefore, the incoming light is reflected back from the metal electrode, and an optical interference takes place. The exact light intensity distribution depends on the optical constants of the materials, the thicknesses of various layers, and the wavelength of the propagating light. The total absorption within the active layer of an MDMO–PPV=PCBM (mixed in 1:4 ratio by weight) has been obtained by optical modeling [70], and its spectral dependence at various active layer thicknesses is shown in Figure 10.9b. Thanks to the large absorption coefficient of the materials, a solar cell using an 80 nm active layer absorbs nearly 60% of the incoming light. Interesting result is that the absorbed fraction of the incoming photons changes only negligibly if the active layer thickness is increased to 120 nm, which points out the importance of interference effects as discussed above.
10.3.3 Low Band Gap Polymers Figure 10.9b shows that by increasing the thickness of the MDMO–PPV–PCBM photoactive layer to 320 nm, 80% of the incoming light can be absorbed up to the wavelength of 600 nm. In comparison to the incoming photon flux of the standard AM (air mass) AM 1.5 illumination, however, it is clear that most of the photons above 600 nm, where the solar photon flux is the most intense, are not utilized [71]. Synthesizing low band gap polymers, therefore, is identified as a key research target [72] toward improved harvesting of the solar spectra, and some encouraging recent examples will be briefly presented. 1. PTPTB is among the first low band gap polymers exhibiting promising photovoltaic characteristics in bulk heterojunction solar cells [73]. It consists of an alternating electron rich N-dodecyl2,5-bis(20 -thienyl)pyrrole and an electron-deficient 2,1,3-benzothiadiazole unit as shown in
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Iin Iout Glass
ITO
(a)
PEDOT -PSS
Active layer
A1
5 3 1018
1.0
4 3 1018
Absorbed fraction
0.8 0.7
20 nm 40 nm 80 nm 120 nm 160 nm 200 nm 320 nm
0.6 0.5 0.4 0.3
3 3 1018
2 3 1018
1 3 1018
0.2
Photon flux (# m−2 s−1 nm−1)
0.9
0.1 0.0 300 (b)
400
500
600
700
0 800
Wavelength l (nm)
FIGURE 10.9 (a) Schematic structure of a multilayer bulk heterojunction solar cell. The incoming light Iin entering from the ITO side is reflected back from the metal electrode. (b) The fraction of absorbed incident light within the MDMO-PPV:PCBM active layer calculated for various thicknesses. The photon flux of a standard AM (air mass) 1.5 solar spectrum is shown for comparison. (Reprinted from Hoppe, H., Arnold, N., Sariciftci, N.S., and Meissner, D., Sol. Ener. Mater. Sol. Cells, 80, 105, 2003. With permission. Copyright 2003. Elsevier.)
Figure 10.10. The onset of the absorption of PTPTB is around 780 nm, which corresponds to an optical band gap of 1.6 eV. When mixed with the electron acceptor PCBM, the characteristic spectroscopic signatures of photoinduced charges are observed [74]. Spectrally resolved photocurrent measurements of a device based on PTPTB=PCBM show clearly that the spectral sensitivity is extended to longer wavelength by almost 200 nm as compared to the MDMO-PPV–PCBM mixture. The power conversion efficiency of 1% is limited by the low FF. 2. Poly(thienylene vinylene)s (PTVs) are interesting class of conjugated polymers that exhibit one of the highest charge mobility in an field effect transistor (FET) structure [75], and typically have a low band gap around 1.6 eV [76]. These two features make them very attractive for photovoltaic applications. Unfortunately, PTVs are insoluble. Vanderzande et al. have developed a novel precursor approach to bulk heterojunction solar cells based on 3,4-dichloro and 3,4-dibromo derivatives of poly(thienylene vinylene)s [77]. A precursor polymer (Figure 10.10) is synthesized via the sulfinyl route, which can be dissolved and mixed with the electron acceptor PCBM in the same solvent. After film deposition by the spin-coating technique, the precursor polymer contained within the mixture is converted to the conjugated polymer PTV by a mild heat treatment. The band gap of the resulting PTV derivative is 1.55 eV. The power conversion efficiency of bulk heterojunction solar cells based on the PTV derivative and PCBM prepared via such precursor route is around 0.2%, and mainly limited by the small open circuit voltage and the low FF. A promising strategy for
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N
S
S N
N
N
R
R S
S
S S
N
n
C12H25
n=1–4
S N
PTPTB
n
N
APFO-Green2
R
R
O
R
S R' S
R
∆T S
n
n
R=Cl or Br Precursor polymer
PTV
FIGURE 10.10 Chemical structure of low bandgap polymers. PTPTB, (poly-N-dodecyl-2,5,-bis(20 -thienyl)pyrrole,2,1,3-benzothiadiazole; PTV, poly(thienylene vinylene).
further improvement of that precursor approach is optimizing the thermal conversion process. It was demonstrated that a postproduction treatment of the diodes can improve the Voc and FF of the diodes, possibly due to reducing shunts and pinholes in the films of the photoactive layer. 3. Photovoltaic devices with spectral response extended to 850 nm have been demonstrated by Zhang et al. based on a fluorine copolymer (APFO-Green2)=PCBM mixture [78]. Although the luminescence of the APFO-Green2 is not completely quenched in the mixture with PCBM, the spectral response of the APFO-Green2=PCBM diodes follows the absorption spectra of the materials. Therefore, the authors concluded that the observed photocurrent is primarily due to photoinduced charge transfer between the components. It is particularly interesting that the open circuit voltage achieved by the APFO-Green2=PCBM device is close to 900 mV, which is among the highest reported for bulk heterojunction solar cells. The rather low FF (0.4) and the short circuit current density around 3 mA cm2 limit the power conversion efficiency to around 1%. The above examples demonstrate that the spectral response of conjugated polymer bulk heterojunction solar cell can be extended up to 850 nm by using low band gap polymers. The improved spectral response, however, does not automatically lead to higher power conversion efficiency of the devices. In bulk heterojunction solar cell, the materials perform multifunctional operation (light absorption, charge generation, and transport of the charges), and in addition they should also be processable from solution. One important function is to transport the photogenerated charges to the electrodes without recombination, which will be discussed in Section 10.4.
10.4
Charge Transport and Recombination
Optimizing the nanomorphology of the phase separation in donor–acceptor blends ensures that all photoexcitations reach a donor–acceptor interface within the exciton lifetime before they relax back to the ground state. The ultrafast timescale of the photoinduced charge transfer at the donor–acceptor interface yields charge generation quantum yield close to unity. Tuning the band gap of the conjugated polymer to lower energies enables efficient photon harvesting of the solar spectrum. The remaining challenge is how to extract the photogenerated charge carriers from the bicontinous, interpenetrating network of electron donor–electron acceptor without significant recombination. Charge collection at the electrodes without losses requires that the transit time ttr of a charge drifting through the interelectrode distance d under the influence of an electric field E is shorter than the lifetime
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of that charge carrier t. The transit time can be calculated from the charge mobility m as ttr ¼ d=Em. The product of mt is an important parameter in solar cells determining the transport distance of the charge as ldrift ¼ mtE or when diffusion is considered, ldiff ¼ (Dt)0.5, where D is the diffusion coefficient and is related to the mobility via the Einstein ratio as D ¼ mkBT=e, where kB is the Boltzmann constant, T is the temperature, and e is the elementary change. In this section, after a brief discussion of the main parameters affecting the transport and recombination of the charges, the results of various techniques that have been used to determine the mobility and lifetime in organic materials are compared. Using the time-of-flight (TOF) and the charge extraction by linearly increasing voltage technique (CELIV) techniques, we show that the recombination in the MDMO–PPV=PCBM bulk heterojunction is nearly Langevin-type, i.e., controlled by diffusion of the charge carriers toward each other. In the highly efficient P3HT–PCBM mixtures, on the other hand, the bimolecular recombination is greatly reduced (non-Langevin-type recombination).
10.4.1 Charge Carrier Mobility 10.4.1.1 Theory and Experimental Techniques As defined by the drift-diffusion equation, the mobility m ¼ etsc=m* is the proportionality constant between the drift current density jd ¼ enmE and the electric field E, where n is the charge carrier concentration and e is the elementary charge. It is ideally a material constant, and depends only on the scattering time between collisions tsc and the effective mass of the electron m*. In heavily disordered organic semiconducting materials, the mean free path of charge carriers is comparable to the intersite distance, thus the charges are localized [79]. Their motion over macroscopic distances is viewed as a sequence of incoherent, thermally assisted hopping events. To calculate the hopping rate between sites i and j with energies «i and «j, the Marcus form of electron transfer (Equation 10.7) should be applicable. For most calculations, a more simple form introduced by Miller–Abrahams (M–A form) in 1960 is used [80]: « « DRij exp jkT i nij ¼ n0 exp 2ga 1; a
«j > «i : «j < «i
(10:8)
The assumption used in the M–A form is that the free energy gain of electron transfer can always be dissipated via the large energy phonons of organic molecules and polymers, and, therefore, hops to sites lower in energy «j < «i are not thermally activated, and will not be influenced by an external electric field. In polycrystalline and amorphous organic materials, disorder attributed to variation in the positions and local orientations of the polymer chains (positional disorder) and variation of the conjugation lengths and differences in the local environment (energetic disorder) splits the energy levels to a manifold of localized states. As a consequence, the hopping rate on the macroscopic scale may not be a constant, but is subject to a distribution function [81]. The direct consequence of molecular-scale disorder (and the distribution of hopping rates) is that the mobility is no longer a material constant, but depends on the electric field, the occupational density (charge concentration), the temperature, and in thermal nonequilibrium conditions it may also depend on the arrival time of the migrating carriers, resulting in a thickness-dependent dispersive transport [82]. It was shown that if the shape of the DOS is a Gaussian, the relaxing photoexcitations may reach quasiequilibrium at longer timescales [83]. Monte Carlo techniques [84], an effective medium theory [85], and more recently an analytical approach [86] have predicted that the temperature and electric field dependence of the equilibrium mobility should follow the form
2 s 2 s 2 2 1=2 exp C S E m(T ,E) ¼ m0 exp 3 kT kT
(10:9)
where S is a parameter characterizing positional disorder, m0 (cm2 V1 s1) is a prefactor mobility in the energetically disorder-free system, E (V cm1) is the electric field, and C is a fit parameter. Several
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extensions to Equation 10.9 have been proposed, including a correlated disorder model [87] or a model that can also account for the concentration dependence of mobility [88]. The consequence of the disorder is that the experimental conditions (nature of charge generation, electric field range, and charge concentration [89]) are expected to influence the determined mobility values even in the same materials. A good example is the orders of magnitude higher charge carrier mobility observed in a FET structure as compared to space charge limited current (SCLC) measurements [90] or mobility measured by the TOF technique [91]. A further complication that prevents the unification of the mobility values measured by FET and other techniques is that the charge mobility may be sensitive to the interface chemistry of the dielectric used [92]. The different orientation of crystallites near the surface of the insulator may also lead to different mobility values parallel or perpendicular to the substrate [62]. The charge mobility determined by various experimental techniques in bulk heterojunction solar cells, and in their individual components are summarized in Table 10.1. Charge mobility in MDMO–PPV and P3HT is one of the most intensively investigated, and only a few selected references have been included for comparison. As clearly seen from this data, the mobility determined for nominally the same materials also differ between various measurements. For example, note the rather large difference in electron mobility of pristine PCBM in a FET structure, or the electron and hole mobility in the MDMO–PPV=PCBM mixture determined by the TOF technique. These differences are often attributed to differences in material purity, and the exact sample preparation conditions, such as the solvent used, evaporation rate of the solvent, and posttreatment procedures. Our strategy is to determine the mobility values using experimental techniques that are closely related to the specific application considered, i.e., the photovoltaic mode of thin film photodiodes. In addition, TABLE 10.1 Electron and Hole Mobility of Materials Commonly Used for Bulk Heterojunction Solar Cells and in Their Photovoltaic Blends Material MDMO-PPV MDMO-PPV MDMO-PPV MDMO-PPV P3HT P3HT P3HT P3HT PCBM PCBM PCBM MDMO-PPV=PCBM MDMO-PPV=PCBM MDMO-PPV=PCBM MDMO-PPV=PCBM P3HT=PCBM P3HT=PCBM P3HT=PCBM
Electron Mobility (cm2 V1 s1)
Hole Mobility (cm2 V1 s1)
2103 N=A N=A N=A N=A 6104 1.5104 N=A 2103 4103 1102 4104 2103 2103
5104 5107 2.4106 8.5106 4.7103 2104 3104 3104 N=A N=A 8103 2106 2104 2104
2104 (charge polarity cannot be distinguished) 8.3102 (sum of mobilities Sm are estimated) 4103 N=A 6105 1104 4 8105 310
Technique Used
Comment
FET [92] SCLC [93] TOF [94] TOF [95] FET [96] FET [92] TOF [97] TOF [91], CELIV [98] SCLC [52] FET [99] FET [100] TOF [94] DI-SCLC, TOF [101] SCLC, impedance, tr-EL [102] Photo-CELIV [104,105]
M-XYL, 2008C N=A, RT, T dep. CB, RT CB, RT, T dep. CB, RT M-XYL, 1008C CHF, RT CHF, RT, T dep. CB, RT CHF, RT CB, RT CB, 33:67, RT, # CB, 33:67, RT, # CB, 20:80, RT, # in Ref [103] CB, 20:80, RT, T dep
FP-TRMC [26]
DCB, 25:75, RT, #
TOF [106] TOF [107] FET [108]
N=A, 33:67, RT N=A, 50:50, RT, # CHF, 65:35, RT, #
FET, field effect transistor; SCLC, space charge limited current; TOF, time-of-flight; CELIV, charge extraction by linearly increasing voltage; FP-TRMC, flash photolysis–time resolved microwave conductivity; DI SCLC, Dark injection space charge limited current; tr-EL, transient electroluminescence; M-XYL, mixed xylenes; CB, chlorobenzene; CHF, chloroform; DCB, 1,2-dichlorobenze; RT, room temperature. The ratio given by x:y is the weight percent of the materials when used in the blend, the symbol # indicates if a weight ratio dependent study is available. A detailed temperature dependence of the mobility study is also indicated when available.
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methods that can be used on operational photovoltaic cells are preferred. In the following, we demonstrate that by using TOF, CELIV, and Photo-CELIV techniques, both the charge carrier mobility and recombination mechanism can be studied simultaneously in operational solar cells. 10.4.1.2 Transient Charge Extraction Techniques: A Comparative Study of TOF and CELIV The experimental setups of TOF and CELIV and their typical electrical response are schematically illustrated in Figure 10.11. In the TOF, the sandwich-type structure is illuminated through a transparent electrode by a short laser flash, and charges are generated within a narrow charge generation region lG. The photogenerated charges drift the interelectrode distance d under the applied external field E, which give rise to a current transient. If the charge generation region is much thinner than the interelectrode distance lG
Time-of-flight
Photocurrent
lG = d/10 t tr
Time Oscilloscope
lG = d
Current ( j )
CELIV ∆j j(0)
Time tmax Oscilloscope
FIGURE 10.11 Experimental setups and the typical response of the time-of-flight and the charge extraction by linearly increasing voltage technique (CELIV).
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[109], is that the dielectric relaxation time ts ¼ ««0=s related to the conductivity s ¼ nem should be larger than the transit time of the charge carriers, otherwise the charges that are extracted when the DC voltage is applied to the sample are sufficient to significantly redistribute the electric field leading to anomalous (sometimes negative) electric field dependence of the mobility. Jusˇka and coworkers have developed a complementary technique to TOF, in which the charges that cause the conductivity in the dark are extracted by a triangular voltage pulse applied to a blocking contact of the sample (CELIV) [110]. The extraction field pulls the charges toward the electrodes, which gives rise to a current transient as schematically illustrated in Figure 10.11. The initial and the end of the pulse current value corresponds to the capacitive displacement current j(0) ¼ A ««0=d if all charge carriers are extracted during the pulse and the generation during extraction is negligible. By selecting the proper voltage rise speed A ¼U=t, therefore, the current goes over a maximum, and from the time to reach the maximum extraction current tmax the mobility can becalculated according to [110]
m¼ 2 3Atmax
2d 2
if Dj 1 þ 0:36 j(0)
Dj j(0)
(10:10)
Equation 10.10 has been derived by solving the continuity, current, and Poisson equations [111]. The expression 1 þ 0.36 Dj=j(0) in the denominator has been introduced to compensate for the redistribution of the electric field during charge extraction, and it is valid for moderately conductive samples, i.e., when the number of extracted charge carriers equals or is less than the capacitive charge Dj j(0). The advantage of CELIV is that samples with small thicknesses can be measured, as the charge generation (typically by chemical doping in organic materials) is uniform throughout the whole thickness (lG ¼ d).* As argued above, the TOF used for low conductivity samples and CELIV for reasonably conductive samples are normally complimentary to each other. The conductivity of regioregular P3HT is known to be sensitive to exposure to air, which is due to doping of P3HT by oxygen or humidity [96,112]. We could, therefore, compare the mobility using the TOF technique in P3HT samples prepared and stored in vacuum, with the mobility determined by the CELIV technique on same samples that have been exposed to air, i.e., slightly doped [98]. This comparison showed that the TOF and CELIV yields mutually consistent picture of mobility in the studied P3HT samples. Importantly, the negative electric field dependence of mobility [91] at higher temperatures is observed using both experimental techniques, which confirms that it is not an artifact of the TOF technique, but rather an intrinsic property of the materials studied.
10.4.2 Charge Mobility and Recombination Studied by the Photo-CELIV Technique 10.4.2.1 Nongeminate Charge Recombination The photogenerated electrons and holes are annihilated by geminate and nongeminate recombinations. Three main processes may control the nongeminate recombination rate: 1. Transport of the charges to the donor–acceptor interface. 2. If a potential barrier is formed at the donor–acceptor interface due to the presence of a dipole layer in the dark, the charge needs to overcome that barrier [24]. 3. The recombination reaction between the oppositely charged particles is a bimolecular electron transfer reaction, and the back electron transfer rate may be described by the Marcus-type electron transfer rate [69]. The kinetic order of the whole recombination process is controlled by the kinetic order of the ratelimiting step. *If the charge generation is not uniform, mobility can be calculated using different form of Equation 10.10.
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Generally, the rate equation of a bimolecular recombination reaction is written as dn dp ¼ ¼ bnp, dt dt
(10:11)
where n and p are the concentration of electrons and holes, respectively, and b is the bimolecular recombination coefficient. If the bimolecular recombination of the charges is controlled by the diffusion of the charge carriers toward each other, then the recombination is often the Langevin-type, and the bimolecular recombination coefficient bL can be written as bL ¼ e(meþmh)=««0, where me and mh are the electron and hole mobilities, respectively. 10.4.2.2 Experimental Arrangement of the Photo-CELIV The schematic response of the photo-CELIV measurement is illustrated in Figure 10.12. Upon the application of a reverse bias (ITO connected to the negative terminal) linearly increasing voltage pulse, the typical current transient is a rectangular-shaped capacitive displacement current with the plateau value j(0) in the dark. When a strongly absorbed short laser flash hits the sample, charges are photogenerated throughout the sample. The charges either undergo recombination or are extracted from the
U
Umax
0V Uoffset t t1/2 j
∆j
j(0)
tdel
0
tmax
t
FIGURE 10.12 Pulse sequence and schematic response of the photo-CELIV technique. (Reprinted from Mozer, ¨ sterbacka, and G. Jusˇka, Phys. Rev. B., 72, 035217, A.J., G. Dennler, N.S., Sariciftci, M., Westerling, A., Pivrikas, R. O 2005. With permission of American Physical Society.)
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200 160 120 80 40 0 −40 −80 −120 −160 −200
U offset 0V
U offset 0.9 V
light voltage −5x10−6
5x10−6
0
−6 1x10−5 −5x10
Time (S)
FIGURE 10.13
U offset 1.2 V
0
−6 5x10−6 1x10−5 −5x10
Time (S)
5x10−6
0
1x10−5
Time (S)
The effect of Uoffset on the photo-CELIV transients recorded for an MDMO-PPV=PCBM solar cell.
device under the influence of the built-in electric field. The built-in field can be compensated by the application of a forward bias offset voltage leading to flat band conditions, and the charges are annihilated by recombination. If a reverse bias, triangular-shaped voltage pulse is applied to the sample after a delay time tdel, but shorter than the charge carrier lifetime, the charge carriers that survived recombination after tdel can be extracted as shown in Figure 10.12. From tmax, the mobility is calculated, and by integrating the photo-CELIV curves after illumination and in the dark, the concentration decay of the charge carriers is monitored. The effect of Uoffset on the transient current response is illustrated in Figure 10.13. In the left hand side at zero offset bias (short circuit), the majority of photogenerated charge carriers exit the devices before the voltage pulse is applied. In the middle graph, the Uoffset was so tuned that the photocurrent immediately upon photoexcitation is minimal (flat band condition). The right hand side shows a transient when the Uoffset is larger than the flat band potential. As a result, charge carriers are injected from the contacts in the dark (note the nonzero current baseline) and the photogenerated and injected charges are extracted to the opposite direction as compared to the photovoltaic mode. The change in the sign of the extraction current within a small voltage range above the flat-band potential clearly indicates that the drift component of the photocurrent dominates the transient response. Photo-CELIV transients at various delay times have been recorded at room temperature as shown in Figure 10.14. The maximum of the extraction pulse Dj is decreasing as tdel is increased, which is related
t del
90
2 µs 4 µs
Extraction current (Am−2)
6 µs 8 µs 14 µs
60
24 µs 40 µs dark 30
0 0
2x10−6
4x10−6
6x10−6
8x10−6
1x10−5
1x10−5
Time (s)
FIGURE 10.14 Photo-CELIV transients recorded in MDMO–PPV=PCBM (1:4) bulk heterojunction solar cells at various delay times at room temperature.
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to charge recombination. The time to reach the maximum extraction current tmax is increasing; therefore, the mobility calculated according to Equation 10.10 is decreasing as the time delay is increased. 10.4.2.3 Time-Dependent Mobility and Recombination in Bulk Heterojunction Solar Cells Similarly to the extraction transients recoded at room temperature, Dj is decreasing and tmax shifts to longer times with increasing delay time at all measured temperatures. Contrary to the room temperature transients, the end of the pulse extraction current at lower temperatures (<150 K) does not reach the capacitive value j(0) indicating that some portion of the charge carriers are trapped during the applied extraction pulse. The mobility values and the concentration of the extracted charge carriers versus delay time are plotted for various temperatures in Figure 10.15. The decreasing mobility with increasing delay time is attributed to energy relaxation of the charge carriers toward the tail states of the distribution [105].
−3.2
300 K, 250 nm 300 K, 265 nm 300 K, 360 nm 270 K, 360 nm 240 K, 360 nm 210 K, 360 nm 180 K, 360 nm 150 K, 360 nm 120 K, 360 nm
log10 m (cm2 V−1 s−1)
−3.6 −4.0 −4.4 −4.8 −5.2 −5.6 −6.0
(a)
300 K, 250 nm 300 K, 265 nm 300 K, 360 nm 270 K, 360 nm 240 K, 360 nm 210 K, 360 nm 180 K, 360 nm 150 K, 360 nm 120 K, 360 nm
n (t) (cm−3)
1016
(b)
1015
1014 10−7
10−6
10−5
10−4
10−3
10−2
10−1
100
101
tdel(s)
FIGURE 10.15 The mobility and concentration of extracted charge carriers versus the delay time at various ¨ sterbacka, temperatures. (Reprinted from Mozer, A.J., G. Dennler, N.S., Sariciftci, M., Westerling, A., Pivrikas, R. O and G. Jusˇka, Phys. Rev. B., 72, 035217, 2005. With permission of American Physical Society.)
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Note that charge carriers can be extracted even after several milliseconds at lower temperatures, which is consistent with the frequently observed millisecond lifetime often observed in transient absorption and photoinduced absorption studies at lower temperatures. The concentration decay in Figure 10.15b at lower temperatures does not follow a simple bimolecular recombination. Instead, Equation 10.12 is used, which, based on the observed time dependence of the mobility, uses a time-dependent bimolecular recombination coefficient b(t) as n(t) ¼ p(t) ¼
n(0) Rt 1 þ n(0) b(t)dt
(10:12)
0
It occurs that the time dependence of b(t) obtained by fitting follows a power law at all temperatures as b(t) ¼ b0 t(1-g), where b0 and g are temperature dependent parameters, the latter related to the magnitude of dispersion. g is close to 1 at room temperature, indicating time-independent bimolecular recombination at this higher temperatures, but gradually decreasing at lower temperatures. It is tempting to directly compare the measured time-dependence of the mobility m(t) with the timedependence of the bimolecular recombination coefficient b(t) using the photo-CELIV technique. In the case of Langevin-type bimolecular charge carrier recombination, b is proportional to the mobility as b(t) ¼ Bm(t), where B (V cm) is a constant B ¼ e=««0 ¼ 6107 V cm (calculated using « ¼ 3 for the MDMO-PPV–PCBM blend). Thus, the concentration decay of charges (Equation 10.12) can be written as a function of the time-dependent mobility: n(t) ¼
n(0) Rt 1 þ n(0)B m(t)dt
(10:13)
0
The fits to the concentration decay at higher temperatures (>210 K) using Equation 10.13 are shown in Figure 10.16, yielding B ¼ 2107 V cm at 300 K, and slightly smaller values when the temperature is decreased. This value is relatively close to the value 6107 V cm, which indicates that Langevin-type bimolecular recombination is operative at room temperature and at the measured microsecond– millisecond timescale. The found diffusion-controlled recombination is consistent with previous results
300 K 270 K 240 K 210 K 180 K
n (t) (cm−3)
1016
1015
10−6
10−5
10−4
10−3
10−2
t (s)
FIGURE 10.16 The concentration decay of the extracted charge carriers fitted using equation 10.13. (Reprinted ¨ sterbacka, and G. Jusˇka, Phys. Rev. B., from Mozer, A.J., G. Dennler, N.S., Sariciftci, M., Westerling, A., Pivrikas, R. O 72, 035217, 2005. With permission of American Physical Society.)
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using transient absorption [113], and transient, nonresonant hole burning spectroscopy [23]. Those results, however, indicated a power law decay of the photogenerated charge carriers on similar timescales rather than the hyperbolic decay of the extracted charges as shown above. The observed difference may be related to the fact that the recombination of all, including the deeply trapped charge carriers is monitored using optical pump-probe techniques, meanwhile only the reasonable mobile charge carriers are extracted in the photo-CELIV technique. We note that the transient absorption studies also identified a faster (<400 ps) recombination, which is largely temperature independent but lightly intensity dependent. Nelson explained that the early time recombination events are due to charges before trapping (above mobility edge), and described the recombination dynamics covering the nanosecond–millisecond timescale using a Monte Carlo technique [114].
10.4.3 Reduced Recombination in the P3HT and PCBM Bulk Heterojunction The implication of a diffusion controlled Langevin-type bimolecular recombination in the MDMO-PPV and PCBM solar cells is that the mobility and the lifetime are coupled to each other. In such cases, increasing the mobility may also enhance the recombination, and the mt product relevant for efficient charge extraction may remain unchanged. In the following, we will show that in the high efficiency P3HT=PCBM solar cell mixtures the bimolecular recombination is drastically reduced as compared to the Langevin recombination, which is good for reaching high power conversion efficiencies [106]. Figure 10.17 shows the TOF photocurrent transients recorded for P3HT and PCBM solar cells with a postproduction treatment following the procedure described in Ref. [56]. The samples were excited by a short, 3 ns laser pulse at 532 nm wavelength, which corresponds to an optical density of 4.4. The RC constant of the setup was larger than the transit time of the charge carriers tRC>>ttr , i.e., the extraction of charge carriers from the sample is limited by the circuit resistance (integral mode TOF). As shown in Figure 10.17, the extracted charge Qe is normalized to the charge stored on the capacitor CU scales linearly with the light intensity at lower light intensities. At higher light intensities, on the other hand, a clear saturation in extracted charge Qe as well as the extraction time te is observed, which is attributed to bimolecular recombination. For these high light intensities, the extracted charge is given as
Qe ¼
Z1 0
je dt ¼
edS , bte
(10:14)
where je is the photocurrent transient, d is the thickness, and S is the electrode area, b is the bimolecular recombination coefficient. Note that if monomolecular recombination limits the extracted charge, te does not saturate with light intensity but has a logarithmic light intensity dependence. From Equation 10.14, the ratio b=bL is directly given as b CU0 ttr ¼ , bL Qe te
(10:15)
and from the saturated values of Qe and te shown in Figure 10.17, b ¼ 21013 cm3 s1 has been calculated. The mobility of the faster carrier m ¼ 4103 cm2 V1 s1 has been determined independently using current-mode TOF (tRC ttr), from which the Langevin bimolecular recombination coefficient value b ¼ 2109 cm3 s1 is calculated. The measured b value is at least four orders of magnitude smaller than the Langevin value, which indicates that the Langevin-type, diffusion-controlled recombination typical for low mobility organic materials is not operable in the high efficiency P3HT=PCBM solar cells. The reduced bimolecular recombination in the P3HT=PCBM mixture has been recently confirmed using a double injection technique [115]. Although the origin of the reduced recombination rate is not
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t (ms) 0.0
0.2
0.4
0.30
L is in arbitrary units L = 100
te
0.25
0.09
L = 3x10−1 L = 10−1
0.20
0.07
L = 3x10−2 L = 10−2
t1/2
0.15
0.05
L = 3x10−3 L = 10−3
0.10
0.04
L = 3x10−4 L = 10−4
0.05
j/jSCLC
j (mA)
0.6 0.11
0.02
0.00
0.00
Qe /CU0
(a) 101
10−1
t1/2 (s)
10−4
10−5 (b), (c)
10−6
10−5
10−4
10−3
10−2
10−1
100
L (arb. units)
FIGURE 10.17 (a) Integral-mode time-of-flight transients recorded for a 700 nm thick, P3HT=PCBM (1:2) solar cells at various light intensities. (b) The extracted charge normalized to the charge stored on the capacitor (c) and the charge extraction time versus normalized light intensity. (Reprinted from Pivrikas, A., G. Jusˇka, A.J. Mozer, ¨ sterbacka, Phys. Rev. Lett., 94, 176806, 2005. M. Scharber, K. Arlausbas, N.S., Sariciftci, H. Stubb, and R. O
completely understood, it is intriguing to attribute it to the specific morphology in the treated P3HT=PCBM blends. Charge carriers reach the P3HT fibrillar nanocrystals present in the annealed films as shown in Figure 10.6 and Figure 10.7 in Section 10. 2, thus they are spatially separated from the negative charge reducing the probability for recombination. Alternatively, it cannot be excluded that the reduced recombination rate is due to an interfacial potential barrier due to a ground state dipole between the electron donor and the acceptor, or may even reflect the intrinsically slow back-electron transfer [69].
10.5
Contacts
The necessary driving force in a solar cell is the gradient in the electrochemical potential that drives the photogenerated charges toward semipermeable membranes [116]. In addition, an electrostatic potential
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may also be present in a thin film of large band gap, intrinsic organic semiconductors inserted between two electrodes with asymmetric work functions (metal-insulator-metal, MIM, structure) [117]. The electric field gives rise to a drift current and may also contribute to charge generation by facilitating exciton dissociation or charge separation of the Coulomb coupled e–h pairs [118]. First, the experimental techniques that probe the internal electric field in thin film organic diodes are presented. A practically more relevant discussion of how the material and contact properties influence the maximum Voc in bulk heterojunction solar cells is also included.
10.5.1 The Metal–Insulator–Metal Structures The injection current in thin conjugated polymer devices with various layer thickness inserted between electrodes with different work-function scales with the applied electric field, suggest that a tunnelingbased electron injection is operating [117]. The turn-on voltage defined as the externally applied voltage leading to flat band conditions, on the other hand, is thickness independent and depends on the charge transfer gap (approximately the HOMO–LUMO gap), and the work functions of the electrodes as shown in Figure 10.18. The turn-on voltage in an LED, therefore, is good measure of the flat band potential, and can also be used to estimate the built-in electric field. Because of the disorder that splits the HOMO– LUMO levels to a distribution of states, the estimation of the flat band potential from the turn on voltage is not always straightforward [119]. The electroabsorption technique introduced by Campbell et al. [120] can be used to estimate the built-in electric field in MIM diodes in a more elegant way. The fundamental electroabsorption response is proportional to the square root of the applied electric field and the imaginary part of the third-order susceptibility Im x. The electro-modulation spectrum resembles the first derivative of absorption. The normalized change in the transmission of the sample with a sinusoidally applied electric field is jDT j (hn) / (VDC VBI ) VAC sin (vt), T
(10:16)
where v is the frequency of the modulated voltage superimposed on the DC offset bias VDC. Therefore, DT vanishes when a step-wise varied VDC superimposed on a constant frequency and amplitude AC voltage cancels the internal field. As an example, the electric field–induced change in the transmission of a thin MDMO–PPV layer sandwiched between an ITO-covered glass and evaporated Al contacts at various applied VDC is shown in Figure 10.19. In good agreement with theoretical treatment, the spectral shape of the electro modulation spectrum matches well with the first derivative of the MDMO–PPV absorption at all applied voltages [121]. From Figure 10.19, the internal electric field of 1.3 V is determined. In other samples, additional features due to the modulation of injected or photogenerated charges, or due to electroluminescence may
Ca Ca ITO
Ca
ITO
1.8eV
ITO
2.1eV
Zero bias
Flat-band condition, the onset of injection
Forward bias, tunneling of both carriers
FIGURE 10.18 The band diagram at zero bias (right), flat-band conditions (middle), and at forward bias (right) illustrated schematically according to the MIM picture for the conjugated polymer MEH-PPV.
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1.0
−∆T/T / 10−4
0.5 0.0 −0.5 −1.0 −1.5 −2.0 500 550 600 Wa
vel
650
eng
th (
nm
)
700 −8
−6
−4
−2
0
2
4
6
V)
t(
e ffs
O
FIGURE 10.19 Electroabsorption spectra of an ITO-MDMO-PPV-Al device at various DC bias levels measured at fixed AC voltage and frequency.
also contribute to the DT spectrum making the determination of the internal electric field, especially near and above the flat band potential where charge injection is more pronounced, problematic [122]. Heller et al. have shown that the maximum built-in potential of MEH–PPV layer sandwiched between a set of metal electrodes with unsymmetrical work functions (Al–Ca, Cu–Ca, Au–Ca and Pt–Ca) is approximately 2.1 V, corresponding roughly to the single particle energy gap of the semiconductor in correspondence with the MIM model. If, however, 5 w % C60 is mixed with the conjugated polymer, the maximum built-in potential is reduced by as much as 0.6 V. Pinning of the chemical potential of the anode electrode via charge transfer from the metal to C60 has been proposed [123]. Interestingly, the determined built-in potential scaled with the work function of the higher work-function electrode and, therefore, the work function of this electrode is not pinned to any of the levels of the semiconductor.
10.5.2 Origin of the Voc in Bulk Heterojunction Solar Cells The maximum open circuit voltage is given by the maximum splitting of the electrochemical energy of the electron–hole pair under illumination, in bulk heterojunction solar cells approximately corresponding to the difference between the LUMO level of the acceptor and the HOMO of the electron donor. This has been confirmed by measuring the Voc of bulk heterojunction solar cells based on electron acceptors with varying first reduction potentials as shown in Figure 10.20a [124]. More recently, Gadisa et al. fabricated bulk heterojunction solar cells based on polythiophenes with various oxidation potentials, and showed that the Voc is increased by decreasing the HOMO level of the conjugated polymer [125]. The important issue that we further consider is how this maximum Voc is influenced by the contact properties, which has been the subject of a systematic investigation by changing the electrode materials and thus varying the work function of both the electron and hole contacts. Figure 10.20b shows the variation of the measured Voc for MDMO–PPV=PCBM bulk heterojunction solar cells using Ca (f ¼ 2.87 eV), Al (f ¼ 4.28 eV), Ag (f ¼ 4.26 eV), and Au (f ¼ 5.1 eV) as top electrode, where f is the work function of the metals. The work function of the metals used changes by more than 2 eV, nevertheless only very slight variation of the Voc is observed. The weak dependence of Voc on the work function of the top electrode cannot be explained by a simple MIM picture. The authors proposed that the work function of the top electrode is pinned to the LUMO level of the PCBM, similar to the paper published by Heller et al. [123]. More recently, Mihailetchi et al. investigated the Voc dependence on the
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0.90
0.85 PCBM
0.85
S1 = 0.95
0.80
0.75 0.70 0.65
C60
Azefulleroid
0.75 0.70
Al Ca Au
S2~0.1
Ag
0.65 0.60
0.60
0.55
Ketolactam
0.55 −0.70 (a)
Voltage (V)
Voltage (V)
0.80
−0.65
−0,60
−0,55
E1Red(V)
0.50 2.5 (b)
3.0
3.5
4.0
4.5
5.0
5.5
Work function (eV)
FIGURE 10.20 (a) The experimentally observed Voc of bulk heterojunction solar cells versus the first reduction potential of various electron acceptors, and (b) the work function of the evaporated top electrode. (Reprinted from Brabec, C.J., Cravino, A., Meissner, D., Saricifcti, N.S., Fromherz, T., Rispens, M.T., Sanchez, L., and Hummelen, J.C., Adv. Func. Mater., 11, 374, 2001. With permission from Wiley–VCH, Weinheim, Germany.).
work function of the top electrode using other metals, such as palladium (Pd), which has a very similar work function as Au [126]. They showed that the Voc in the device using Pd as top electrode is reduced to 0.3 eV, and argued that a MIM picture can explain the results. Frohne et al. varied the work function of the hole contact Pedot –PSS by ex situ adjusting its work function by controlled chemical doping [127]. They found that the Voc directly scales with the work function difference of the PEDOT–PSS and the Al contact used. Interestingly, the Voc reversed sign when the work function of the PEDOT–PSS was lower than the LUMO of PCBM, which further suggests the pinning of the top electrode work function. It also demonstrates that the hole contacts is not pinned to any of the electronic levels of the conjugated polymer or PCBM. To summarize, the following empirical expressions for the Voc have been proposed [128]: Voc ¼ fM1 E red(A) only if Aox > fM1 Voc ¼ Aox E red(A) only if Aox < fM1
(10:17)
where Aox is the oxidation potential of the electron donor, Ered is the reduction potential of the electron acceptor, and fM1 is the work function of the hole contact. Finally, it has been demonstrated that insertion of thin (<15 nm) LiF layer below the metal top contact increases the FF of the solar cells considerably [129]. The increased performance is controversial. First, lowering series resistance by forming a better ohmic contact to the LUMO of the PCBM has been proposed and discussed. Furthermore, a light doping of Li can influence the first monolayers below electrode. Finally, dipole layers formed at the interface can reinforce or counteract the charge carrier extraction, thereby manipulating the contact efficiency. As the used LiF layer thickness is below a closed layer minimum thickness, there is high probability that only islands are formed and that such islands bring up an improved contact is still puzzling. Nevertheless, phenomenologically the LiF=Al contacts, therefore, provide a good contact achieving high FFs and high injection current in polymer solar cells [129], polymer LEDs [130] as well as in PCBM-based organic diodes [131].
10.6
Device Models of Bulk Heterojunction Solar Cells
10.6.1 The Equivalent Circuit Model All the important processes of light absorption, charge transport, recombination and contact properties discussed previously can be incorporated into a device model. Such a model is highly desired as the
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performance-limiting factors could be more readily identified, and the theoretical limits of currently used materials estimated. Nevertheless, this area of research as it will be briefly shown here, remains probably the most controversial within the device physics of bulk heterojunction solar cells. We begin the discussion with a simple one-diode equivalent circuit model, which, due to its simplicity and relatively simple procedures to fit current–voltage curves of various polymer solar cells has been frequently applied. The current density–voltage curves of MDMO-PPV:PCBM (1:4) bulk heterojunction solar cells based on the commonly used regiorandom RRa-MDMO-PPV and a novel regiospecific RS-MDMO-PPV synthesized by the group of Vanderzande [95] are displayed in Figure 10.21. The power conversion efficiency calculated using Equation 10.18 is around 2.5% for the RRa-MDMO-PPV:PCBM device, and slightly higher for the RS-MDMO-PPV device, because of the higher (0.71) FF: hAM1:5 ¼
jsc Voc FF m, Pin
(10:18)
where jsc and Voc are the short circuit current density and open circuit voltage, respectively, m is a mismatch factor that accounts for deviations in the spectral response to that of a reference cell. The filling factor, FF, is defined as FF ¼
jMPP VMPP jsc Voc
(10:19)
A simple replacement circuit based on one-diode model of a solar cell, shown in the inset of Figure 10.21, consists of a diode (represented by its quality factor n, and the reverse bias, dark saturation current j0), and a series resistor Rs and a parallel resistor Rp that account for ohmic losses and shunt leakage through the diode, respectively. The photocurrent generation is represented by a current source jph.
0 Voc Rs
Current density (mA cm−2)
−1 V
−2
jph
Rp
n; j0
−3
−4 jsc −5 RS-MDMO-PPV:PCBM (1:4) RRa-MDMO-PPV:PCBM (1:4) −6 0
0.2
0.4
0.6
0.8
Voltage (V)
FIGURE 10.21 Current density–voltage curves of bulk heterojunction solar cells under 80 mW cm2 stimulated AM 1.5 illumination. The inset shows the replacement circuit of a one-diode model.
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TABLE 10.2 Photovoltaic Performance of the Bulk Heterojunction Solar Cells, and the Parameters of Equation 10.20 Obtained by Numerical Calculation
1 2
Jsc mA cm2
Voc V
FF
5.0 5.25
0.8 0.82
0.71 0.61
hAM 1.5 %a
J0 mA cm2
Rs V cm2
Rp V cm2
N
2.65 2.5
6107 6107
1.3 3
2150 950
1.9 2
Source: Reprinted from Mozer, A.J., Denk, P., Scharber, M.C., Nengebaner, H., Sariciftci, N.S., Wagner, P., Lutsen, L., and Vanderzande, D., J. Phys. Chem. B., 108, 5235, 2004. Note: Sample 1 and 2 stands for RS MDMO-PPV and RRa-MDMO-PPV, respectively. a Calculated using Equation 10.20, m ¼ 0.751.
The current through that replacement circuit is given by n e o V jR s (V jRs ) 1 þ jph j ¼ j0 exp nkT Rp
(10:20)
The j–V curves shown in Figure 10.21 have been fitted using the one-diode equation, and the obtained parameters are compared in Table 10.2. The improved FF of the RS-MDMO-PPV based device can be attributed to the lower internal series resistance and the improved shunt resistance. The lower series resistance is proposed to originate from the factor of two higher charge carrier mobility of the RS-MDMO-PPV as compared to the RRaMDMO-PPV at room temperature, which has been determined using the TOF technique [95]. The one-diode replacement circuit model has been also used to extract the diode parameters in bulk heterojunction solar cells with or without the insertion of a thin (<15 nm) evaporated LiF layer, and to present a meaningful analysis pointing out the importance of reducing the contact resistance by LiF [129]. Dyakonov has published several papers [51,132,133] using the equivalent circuit model to analyze the temperature dependence of the current–voltage curves. Nevertheless, Schilinsky et al. [134] and Koster et al. [135] have independently demonstrated that such a simple one-diode model cannot properly account for several other characteristics of bulk heterojunction solar cells, such as the observed light intensity dependence of current–voltage curves or the temperature dependence of the Voc. In the former case, a continuous adjustment of the reverse bias saturation current j0 is necessary to fit the measured current–voltage curves of P3HT and PCBM solar cells at various light intensities. As argued by Schilinsky et al., adjusting the dark parameters of the diode with increasing light intensity cannot be justified. In a more recent case, Koster et al. showed that the temperature dependence of the Voc follows a Voc / kT=e dependence rather than the Voc / nkT=e dependence predicted by Equation 10.20. Both of the above groups argued that a more meaningful fit of their data was obtained by assuming that the photocurrent jph is not constant but depends on the externally applied voltage. However, different origins for the voltage-dependent photocurrent have been proposed and will be reviewed below.
10.6.2 Extended One-Diode Model Schilinsky et al. argued that the photocurrent in bulk heterojunction solar cells is primarily electric-field driven. The electric field in the device changes with the externally applied voltage and can be calculated as E ¼ (VVbi)=d, where Vbi is the built-in electric field. Therefore, the mean drift distance of the charge carriers given by ldrift ¼ mtE also changes with the applied voltage. Full extraction is reached when the drift distance of the charges is d, otherwise only charges generated within the drift distance near the contacts contribute to the external photocurrent. The photogenerated charges which do not reach the electrodes within their lifetime, on the other hand, contribute to photoconductivity through their mobility, which increases the shunt resistance under illumination (photoshunt). The following equation has been introduced to describe the light-intensity dependence of the shunt resistance by Waldauf et al. [136]: Rp ¼
cp Ane em 1 , þ Rpdark d
(10:21)
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where the first term is the photoconductivity weighted by the probability that a charge carrier can penetrate the potential barrier in the wrong direction (opposite to the direction of charge extraction), and termed as contact permeability cp. The simulation of their diodes predicted that the power conversion efficiency could be improved up to 5% by increasing the contact permeability c p, thus reducing the photoshunt.
10.6.3 Electric Field–Dependent Dissociation of the Coulomb-Coupled E–H Pairs Koster et al. [135] and later Mihailetchi et al. [118] argued that because of the small dielectric constant of the materials used in bulk heterojunction solar cells, the photogenerated charges should be bound by their mutual coulomb potential (coulomb-coupled pairs) at the donor–acceptor interface with a binding energy of a few tens of an electron volt. They proposed that the electric field present in the devices (as well as the surrounding thermal bath) facilitates the separation of the Coulomb-coupled radical pair leading to free carriers following the Onsager model. In contrary to the field-dependent extraction presented above, it is assumed that photocurrent in bulk heterojunction solar cells is dominated by the field- and temperature-dependent dissociation of the Coulomb pairs. The current–voltage curves therefore, are described by three regimes: 1. At small voltages below the compensation voltage (V0V < 0.1 V) (V0 is the applied voltage at which the photocurrent jph ¼ jljd is nulled), a linear current versus voltage regime is observed, and attributed mainly to a diffusion-controlled current. 2. As the voltage is increased, yet still at moderate voltages, the current is dominated by the dissociation efficiency of the coupled bound pairs into free carriers. 3. At large negative voltages, the photocurrent represents the full separation and, therefore, the maximum quantum yield. The ratio of the short circuit current to the fully saturated current, as argued, gives directly the ratio of the number of coulomb-pairs versus free carriers at short circuit, which is estimated to be only 60% in the studied MDMO-PPV=PCBM samples. More recently, the model has been expanded to describe the performance of bulk heterojunction solar cells depending on the PCBM concentration. In these recent models, the mobility of electrons and holes determined in independent experiments has been incorporated [103].
10.6.4 Numerical Solution to the Drift-Diffusion Equations Gommans et al. used a complete numerical solution of the drift-diffusion equation to describe the current–voltage curves of bulk heterojunction solar cells at various temperatures [137]. They examined the effect of dissimilar electron and hole mobilities (values taken from the literature) and assumed that the recombination is the Langevin-type bimolecular recombination. Importantly, they assumed a fieldindependent generation rate of the charges upon illumination, and were able to fit the j – V curves reasonably well. They also commented that incorporation of a field-dependent generation (as well recombination) according to the Onsager formula, overestimates the field dependence of the measured photocurrent. Furthermore, the j – V curves were modeled at several temperatures by changing the electron and hole mobility, yet keeping the charge generation rate constant. It was concluded that the field dependent photocurrent in bulk heterojunction solar cells can be entirely attributed to a fielddependent bimolecular recombination, which contradicts to the assumptions used in Refs.[118,135] (field-dependent dissociation of the coulomb-coupled pairs).
10.7
Concluding Remarks and Outlook
In the last five years, there has been an enormous increase in the understanding and performance of polymer–fullerene bulk-heterojunction solar cells. Comprehensive studies have been performed to learn more about the crucial material parameters such as nanomorphology, energy levels, photoexcitations,
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charge transport, and electrode materials. To date, power conversion efficiencies close to 4% are routinely obtained and some laboratories have reported power conversion efficiencies of 5%. Important for the future is new materials and device architectures, aiming at increasing the efficiency to 8%–10%. By combining synthesis, processing, and materials science with device physics and fabrication, there is little doubt that these appealing levels of performance will be achieved in the near future.
Acknowledgment A.J. Mozer thanks Harald Hoppe, Joachim Loos, and Christoph Lungenschmied for providing figures for this work. Their help in the preparation of this manuscript is highly appreciated. A.J. Mozer thanks Gilles Dennler and Christoph Waldauf for discussions and acknowledges the financial support of the Japanese Society for the Promotion of Science.
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58. Kim, Y., S.A. Choulis, J. Nelson, D.D.C. Bradley, S. Cook, and J.R. Durrant. 2005. Device annealing effect in organic solar cells with blends of regioregular poly(3-hexylthiophene) and soluble fullerene. Appl Phys Lett 86:063502. 59. Reyes-Reyes, M., K. Kim, and D.L. Caroll. 2005. High-efficiency photovoltaic devices based on annealed poly(3-hexylthiophene) and 1-(3-methoxycarbonyl)-propyl-1-phenyl-(6,6)C61 blends. Appl Phys Lett 87:083506. 60. Ma, W., C. Yang, X. Gong, K. Lee, and A.J. Heeger. 2005. Thermally stable, efficient polymer solar cells with nanoscale control of the interpenetrating network morphology. Adv Funct Mater 15:1617. ¨ sterbacka, R., C.P. An, X.M. Jiang, and Z.V. Vardeny. 2000. Two-dimensional electronic excita61. O tions in self-assembled conjugated nanocrystals. Science 287:839. 62. Sirringhaus, H., P.J. Brown, R.H. Friend, M.M. Nielsen, K. Bechgaard, B.M.W. Langeveld-Voss, A.J.H. Spiering, R.A.J. Janssen, E.W. Meijer, P. Herwig, and D.M. de Leeuw. 1999. Two-dimensional charge transport in self-organized, high-mobility conjugated polymers. Nature 401:685. 63. Schuller, S., P. Schilinsky, J. Hauch, and C.J. Brabec. 2004. Determination of the degradation constant of bulk heterojunction solar cells by accelerated lifetime measurements. Appl Phys A 79:37. 64. Dennler, G., C. Lungenschmied, H. Neugebauer, N.S. Sariciftci, M. Latre`che, G. Czeremuszkin, and M.R. Wertheimer. A new encapsulation solution for flexible organic solar cells. Thin Sol Films 511–512:349. 65. Wu¨rfel, P. 2005. Physics of solar cells–From principles to new concepts, 140. Weinheim: Wiley-VCH Verlag GmbH & Co. KGaA. 66. Schokley, W., and H.J. Queisser. 1961. Detailed balance limit of efficiency of p-n junction solar cells. J Appl Phys 32:510. 67. Wu¨rfel, P. 2005. Physics of solar cells–From principles to new concepts, 138. Weinheim: Wiley-VCH Verlag GmbH & Co. KGaA. 68. Marcus, R.A. 1993. Electron transfer reactions in chemistry. Theory and experiment. Rev Mod Phys 65:599. 69. Lemaur, V., M. Steel, D. Beljonne, J.-L. Bredas, and J. Cornil. 2005. Photoinduced charge generation and recombination dynamics in model donor=acceptor pairs for organic solar cell applications: A full quantum-chemical treatment. J Am Chem Soc 127:6077. 70. Hoppe, H., N. Arnold, N.S. Sariciftci, and D. Meissner. 2003. Modeling the optical absorption within conjugated polymer=fullerene-based bulk-heterojunction organic solar cells. Sol Ener Mater Sol Cells 80:105. 71. Hoppe, H., and N.S. Sariciftci. 2004. Organic solar cells: An overview. J Mater Res 19:1924. 72. Winder, C., and N.S. Sariciftci. 2004. Low bandgap polymers for photon harvesting in bulk heterojunction solar cells. J Mater Chem 14:1077. 73. Anantharaman, D., J.K.J. van Duren, P.A. van Hal, J.L.J. van Dogen, and R.A.J. Janssen. 2001. Synthesis and characterization of a low bandgap conjugated polymer for bulk heterojunction photovoltaic cells. Adv Funct Mater 11:255. 74. Brabec, C.J., C. Winder, N.S. Sariciftci, J.C. Hummelen, A. Dhanabalan, P.A. van Hal, and R.A.J. Janssen. 2002. A low-bandgap semiconducting polymer for photovoltaic devices and infrared emitting diodes. Adv Funct Mater 12:709. 75. Fuchigami, H., A. Tsumura, and H. Koezuka. 1993. Polythienylenevinylene thin-film transistor with high carrier mobility. Appl Phys Lett 63:1372. 76. Murray, M.M., and A.B. Holmes. 2000. Poly(arylene vinylene)s–synthesis and applications in semicondicting devices Semiconducting polymers: Chemistry, physics and engineering, eds. G. Hadziioannou and P.F. van Hutten. Weinheim: Wiley-VCH. 77. Henckens, A., M. Knipper, I. Polec, J. Manca, L. Lutsen, and D. Vanderzande. 2004. Poly(thienylene vinylene) derivatives as low band gap polymers for photovoltaic applications. Thin Sol Films 451–452:572.
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99. Waldauf, C., P. Schilinsky, M. Perisutti, J. Hauch, and C.J. Brabec. 2003. Solution-processed organic n-type thin-film transistors. Adv Mater 15:2084. 100. Anthopoulos, T.D., C. Tanase, S. Setayesh, E.J. Meijer, J.C. Hummelen, P.W.M. Blom, and D.M. de Leeuw. 2004. Ambipolar organic field-effect transistors based on a solution-processed methanofullerene, Adv Mater 16:2174. 101. Tuladhar, S.M., D. Poplavskyy, S.A. Choulis, J.R. Durrant, D.D.C. Bradley, and J. Nelson. Ambipolar charge transport in films of methanofullerene and poly(phenylenevinylene)=methanofullerene blends. Adv Func Mater 15:1171. 102. Melzer, C., E.J. Koop, V.D. Mihailetchi, and P.W.M. Blom. 2004. Hole transport in poly(phenylene vinylene)=Methanofullerene bulk-heterojunction solar cells. Adv Funct Mater 14:865. 103. Mihailetchi, V.D., L.J.A. Koster, P.W.M. Blom, C. Melzer, B. de Boer, J.K.J. van Duren, and R.A.J. Janssen. 2005. Compositional dependence of the performance of poly(p-phenylene vinylene):methanofullerene bulk-heterojunction solar cells. Adv Funct Mater 15:795. ¨ sterbacka, M. Westerling, and G. Jusˇka. 104. Mozer, A.J., N.S. Sariciftci, L. Lutsen, D. Vanderzande, R. O 2005. Charge transport and recombination in bulk heterojunction solar cells studied by the photoinduced charge extraction in linearly increasing voltage technique. Appl Phys Lett 86:112104. ¨ sterbacka, and G. Jusˇka. 105. Mozer, A.J., G. Dennler, N.S. Sariciftci, M. Westerling, A. Pivrikas, R. O 2005. Time-dependent mobility and recombination of the photoinduced charge carriers in conjugated polymer=fullerene bulk heterojunction solar cells. Phys Rev B 72:035217. 106. Pivrikas, A., G. Jusˇka, A.J. Mozer, M.C. Scharber, K. Arlauskas, N.S. Sariciftci, H. Stubb, and ¨ sterbacka. 2005. Bimolecular recombination coefficient as a sensitive testing parameter for R. O low-mobility solar-cell materials. Phys Rev Lett 94:176806. 107. Huang, J., G. Li, and Y. Yang. 2005. Influence of composition and heat-treatment on the charge transport properties of poly(3-hexyltiophene) and [6,6]-phenyl C61-butyric acid methyl ester blends. Appl Phys Lett 87:112105. 108. Nakamura, J., K. Murata, and K. Takahashi. 2005. Relation between carrier mobility and cell performance in bulk heterojunction solar cells consisting of soluble polythiophene and fullerene derivatives. Appl Phys Lett 87:132105. ¨ sterbacka, and H. Stubb. 2002. Charge transport at low 109. Jusˇka, G., K. Genevicˇius, K. Arlauskas, R. O electric fields in p-conjugated polymers. Phys Rev B 65:233208. ¨ sterbacka, and H. Stubb. 2000. Charge 110. Jusˇka, G., K. Arlauskas, M. Vili unas, K. Genevicˇius, R. O transport in p-conjugated polymers from extraction current transients. Phys Rev B 62:R16235. 111. Jusˇka, G., K. Arlauskas, M. Vili unas, and J. Kocˇka. 2000. Extraction current transients: New method of study of charge transport in microcrystalline silicon. Phys Rev Lett 84:4946. 112. Hoshino, S., M. Yoshida, S. Uemura, T. Kodzasa, N. Takada, T. Kamata, and K. Yase. 2004. Influence of moisture on device characteristics of polythiophene-based field-effect transistors. J Appl Phys 95:5088. 113. Montanari, I., A.F. Nogueira, J. Nelson, J.R. Durrant, C. Winder, M.A. Loi, N.S. Sariciftci, and C.J. Brabec. 2002. Transient optical studies of charge recombination dynamics in a polymer=fullerene composite at room temperature. Appl Phys Lett 81:3001. 114. Nelson, J. 2003. Diffusion-limited recombination in polymer-fullerene blends and its influence on photocurrent collection. Phys Rev B 67:155209. 115. Jusˇka, G., K. Arlauskas, G. Sliauzys, A. Pivrikas, A.J. Mozer, N.S. Sariciftci, M. Scharber, ¨ sterbacka. 2005. Double injection as a technique to study charge carrier transport and R. O recombination in bulk-heterojunction solar-cells. Appl Phys Lett 87:222110. 116. Wu¨rfel, P. 2005. Physics of solar cells–From principles to new concepts, 133. Weinheim: Wiley-VCH Verlag GmbH & Co. KGaA. 117. Parker, I.D. 1994. Carrier tunneling and device characteristics in polymer light-emitting diodes. J Appl Phys 75:1656. 118. Mihailetchi, V.D., L.J.A. Koster, J.C. Hummelen, and P.W.M. Blom. 2004. Photocurrent generation in polymer-fullerene bulk heterojunctions. Phys Rev Lett 93:216601.
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11 Biomedical Applications of Inherently Conducting Polymers (ICPs) 11.1 11.2
Introduction..................................................................... 11-1 Synthesis, Processing, and Fabrication of ICPs ............ 11-2 Biocomposites . Review of Electronic and Redox Properties . Sterilization
11.3
Biomolecular Sensing...................................................... 11-9
11.4 11.5 11.6 11.7 11.8
Biomolecular Actuators ................................................ 11-14 ICPs in Tissue Engineering........................................... 11-15 ICPs Used in Nerve Cell Regeneration........................ 11-18 Biodegradability and Stability ...................................... 11-19 Biomechanical Sensing.................................................. 11-20
Oligonucleotides . Enzymes . Antibodies and Antigens
Biomechanical Actuators: Artificial Muscles and Microactuators . Wearable Prosthetics
P.C. Innis, S.E. Moulton, and G.G. Wallace
11.1
11.9
Conclusions and Future Developments: Bionics and Beyond....................................................... 11-24
Introduction
The identification and development of materials that find application in biomedical applications has a profound effect on our quality of life. It is breakthroughs in this area of materials research that has seen the development of highly effective stents [1,2], bone replacements [3], pacemakers [4], bionic ears [5], and wearable prosthetics [6]. In each of the above examples the material requirements will differ. However, in all cases they must be compatible with the biological environment in which they are to operate. This compatibility may involve molecular and cellular interactions, be at the skeletal level wherein aspects related to wearability, comfort, lightweight, and esthetics become important. The quest to more effectively monitor and manipulate the biosystems requires the creation of better interfaces between the biological and electronic domains. Materials that bridge this interface are finding utility in the emerging field of bionics. Inherently conducting polymers (ICPs) have been shown to be excellent bionic materials providing utility from the biomolecular to the biomechanical level. At the biomolecular level, conducting polymers can be produced in ways that make their integration into implants for tissue engineering (TE) or nerve regeneration possible. At the biomechanical (skeletal) 11-1
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level, fabrication protocols that enable wearable structures containing conducting polymers that function as sensors or mechanical actuators (artificial muscles) have been developed. With biomedical applications in mind, this chapter reviews the important elements of the synthesis and processing of conducting polymers as well as their fabrication into devices. The key properties that make the use of ICPs in biomedical applications an attractive proposition are their electronic and electrochemical switching properties. These important features will be discussed with specific emphasis upon their use as sensors or as actuators from the biomolecular to the biomechanical levels.
11.2
Synthesis, Processing, and Fabrication of ICPs
Those interested in the biomedical applications of ICPs can draw on the numerous advances in synthesis and processing to produce organic conductors with appropriate electronic and mechanical properties and yet have biofunctionality effectively integrated. Conducting polymers such as polypyrroles (PPy) are synthesized according to Figure 11.1. The fact that this polymerization process can be induced in aqueous environments (for pyrrole at least) at low oxidation potentials and at neutral pH values is significant. Such conditions are compatible with maintaining the integrity of biological entities from the molecular to the cellular level. The choice of dopant (A) is important in determining the chemical and biochemical properties [7] as well as the electronic [8] and mechanical [8] properties. The simplest approach to the integration of biomolecules into ICPs involves their use as dopant molecules. At the molecular level, a range of bioactive entities such as nonapeptides [9], antibodies [10], enzymes [11], and even entire living cells [12,13] have been incorporated directly into conducting polymers during electrosynthesis. Although conducting polymers may also be produced via chemical oxidation of the corresponding monomer, this approach does not normally provide the degree of control over the oxidation potential, and therefore the rate of polymerization needed to ensure the integrity of the bioagent. Incorporation of biologically relevant functionality into the ICP may be sensitive to the oxidizing power of the chemical reagent driving polymerization. Electrochemical synthesis provides a direct method of control over the oxidation potential and hence the bioactivity of many incorporated species may be retained. To ensure that accurate control of the potential is maintained a three-electrode potentiostatic approach is normally employed. Using this approach, biological polyelectrolytes such as heparin [14] or hyaluronic acid (HA) [15] have also been incorporated. An interesting feature of conducting polymers containing polyelectrolytes as dopants are that they are electronically conductive while a high percentage of water (as much as 90% w=w) is incorporated into the resultant polymer making them an organic electronic hydrogel. Such gel structures may also be realized by growing the conducting polymer through a preformed gel (Figure 11.2). The support structure of the original gel is retained with water contents changing by just a few percent after integration of the conducting polymer [16,17]. Processable (ethanol soluble) hydrogels based on block copolymers of polyethylene oxide and poly(e-caprolactone) or heat-extrudable gels based on cross-linked polyacrylic acid have also been used as the host matrix for ICPs [18]. Although the above approaches are by far the simplest to achieve experimentally, the use of covalent attachment provides a further degree of chemical stability and the ability to localize the biomolecule on
+
A−
A− N H
Oxidise n = determines degree of doping m = determines molecular weight
FIGURE 11.1
Polypyrrole reaction scheme.
N H
n
... (1) m
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a
FIGURE 11.2
Conducting polymer grown through a preformed polyacrylamide hydrogel.
the polymer surface. A number of strategies that enable covalent attachment of biomolecules have been developed. De Giglio et al. [19] covalently attached polypeptides containing a cysteine residue to PPy as shown in Figure 11.3, effectively activating the surface for efficient culturing of osteoblast cell lines [20]. Covalent attachment of enzymes, as shown in Figure 11.4, has been achieved by Yon-Hin et al. [21]. Antigen molecules [22] have been covalently attached to either polypyrrole or polythiophene backbones by direct substitution of the leaving group N-hydroxyphthalimide, as shown below in Figure 11.5. Antibody fragments [23] have been incorporated using reactive groups at a chemically modified electrode surface to enable coupling via carbodiimide chemistries (Figure 11.6). Biological polyelectrolytes such as heparin [24,25] or HA have also been covalently attached to PPy. In all cases, both stability and localization have been achieved. However, covalent attachment also has an adverse effect—decreasing the electronic conductivity of the conjugated polymer. In addition, covalent attachment often involves the use of a linker molecule—decreasing the intimate association of the biomolecule with the conducting polymer backbone. This can diminish the ability to use changes in electronic properties to monitor biomolecule interactions or use electrical stimuli to control such interactions.
cys −−
SH
++
H
H
H
N
N
+
N
N
N
H
H
H
N
−
S−++ −
cys
FIGURE 11.3 Covalent attachment of cysteine residues to a PPy backbone. (From De Giglio, E., Sabbatini, L., and Zambonin, P.G., J Biomater Sci-Polym Ed 10, 845, 2003.)
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+ DEC, pH 5.5
LiAIH4
N
N
CH2
(CH2)3
CH2
NH2
N O GOX
(CH2)3
C
NH
OH
CO
CN
GOX
FIGURE 11.4 Covalent attachment of glucose oxidase (GOx) to pyrrole. DEC 1-[3-(Dimethylamino)propyl]-3 ethyl carbodiimide hydrochloride.
11.2.1 Biocomposites An alternative approach involves blending of conducting polymers with biomaterials of interest to form biocomposites. This, of course, ideally involves the ICP and biomolecule to be soluble in a compatible solvent that does not denature the biomolecule. Such examples are limited and usually involve polyanilines [26], as they are more postsynthesis processable. ICP polyanilines have, for example, been blended with collagen. Blending of poly(o-ethoxyaniline) and collagen resulted in formation of flexible free-standing semiconducting materials. Polyaniline has also been blended with pluronic acids—poly(ethylene oxide)–poly(propylene oxide)– poly(ethylene oxide) triblock copolymers [27], which have been used in a range of medical and pharmaceutical products. As with all composite materials containing two discrete phases, the issue of phase segregation over time needs to be considered. Another approach to the formation of conductive biocomposites involves chemical oxidation of pyrrole [28] or sulfonated pyrrole [29] in the presence of the biomolecule. Polypyrrole chitosan composites have been formed via chemical oxidation of pyrrole in the presence of chitosan [30]. Chitosan is another commonly used biomaterial to which polyaniline has been chemically grafted to produce materials with conductivities as high as 102 S cm1 (Figure 11.7) [31]. For most biological studies, the ICPs have been prepared in the form of a flat film although some tubular structures have also been prepared. An evolving area of interest is the synthesis of high surface area bioactive colloidal systems [32]. PPy-polystyrene core latex particles with an inner core 600 nm in diameter have been functionalized with N-succinimidyl ester [33,34] and N-hydroxysuccinimide [35] to introduce protein-binding capacity. Functionalization of these particles exhibited a high degree of
O
O O
OH
N
OH
O
x
O
6
O
O
N H
N O
N H
y
O
S
O
O
S
n
FIGURE 11.5 Leaving groups to facilitate attachment of antigen to a PPy copolymer or a polythiophene backbone. (From Zhou, D., Too, C.O., and Wallace, G.G., React Funct Poly, 39, 19, 1999; Collier, J.H., Camp, J.P., Hudson, T.W., Schmidt, J Biomed Mater Res, 50, 574, 2000.)
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N
N
Biomedical Applications of Inherently Conducting Polymers (ICPs)
OH
S
S
S
NH O−C
O−C
Fab’
S
Fab’ S
NH O−C
NH O−C
Polymer Step 1
FIGURE 11.6
Step 2
Step 4
Step 3
The steps involved in the attachment of the antibody fragment to a PPy backbone.
reactivity to amine and thiol groups found in proteins. Exposure of these particles to biotin resulted in effective binding and this subsequently resulted in a system capable of specific binding to avidin protein after incubation or human serum albumin with binding up to 0.2 mg=m2. The direct incorporation of bioactive proteins during electrochemical production of PPy colloids using a novel flow through cell has been demonstrated [36]. Both human serum albumin and a-lactalbumin have been incorporated and have shown to retain bioactivity [37]. In transiting from the biomolecular to the biomechanical (skeletal level) world, the synthesis and fabrication of ICPs as wearable materials becomes the challenge. In terms of wearables, the conducting polymer may be integrated into a preformed textile structure at the molecular level (akin to dying the textile) or preformed ICP fibers can be weaved into a textile structure. The former can be achieved using the in situ polymerization process, originally introduced by Kuhn and coworkers at Milliken [38], as depicted in Figure 11.8. Dopant anions within the ICP coated onto a textile surface can play a dual role in that it can assist in substrate attachment and simultaneously provide molecular level doping to ensure electronic conductivity.
CH2OH
CH2OH
O
2−
O
O
S2O8 / HC1 O
O OH
OH
O HN
NH2
C
NH2 CH3
Chitosan
CH2OH
CH2OH
O
Films cast from aqueous solution
O
O O
O OH
OH
Fibres formed by injection into base
O NH
HN
N H
C
N H
CH3 N H
n
Chitaline
FIGURE 11.7
Partial structure of chitosan and reaction scheme for formation of chitaline graft copolymer.
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Fabric
Oxidant
NDSA
NDSA
M
NDSA
M
M
NDSA
Ox
M
FIGURE 11.8 2005.)
In situ polymerization on textile. (Courtesy Dr Jian Wu, PhD diss., University of Wollongong,
This approach has been shown to have minimal effect on the mechanical properties on a range of base textiles including nylons knitted with Spandex (a registered trademark of E.I. du Pont de Nemours and Co., Wilmington, Delaware). Uniform, adherent conducting and electroactive polymer coatings can be achieved in this way. The redox activity (discussed later) of the polymer-coated textiles is clearly evident in the cyclic voltammogram (Figure 11.9). The underlying properties of the textiles can be used to advantage in that they can induce a biomechanical transduction method as described by De Rossi [39] and used by us in the development
220
Ep(a)
170
Current (mA)
120 70 20 −30 −80
−130 −180 −230 −120
Ep(c) −0.70
−0.20
0.30
0.80
1.30
Potential (V)
FIGURE 11.9 Cyclic voltammogram of Ppy- NDSA-coated Nylon Lycra in 1.0 M NaNO3 at a scan rate of 100 mV s1. The PPy coated Nylon Lycra was prepared using the in situ method for 2 h in an aqueous solution containing 0.015 M pyrrole, 0.04 M FeCl3, and 0.005 M NDSA. (Courtesy Dr Jian Wu PhD diss., University of Wollongong, 2005.)
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11-7
Step 1
Thermometer
Textile is wrapped around spindle
PMAS solution
NH2
Step 2 Add Aniline
Thermometer
Textile is wrapped around spindle
PMAS/aniline solution
Step 3
Add Oxidant
Thermometer
PMAS/aniline solution
FIGURE 11.10
Textile is wrapped around spindle
Molecular templating on textile.
of a biomechanical feedback device [40]. Alternatively, textile substrates known to be biodegradable (such as polyesters) can be used as a platform for conducting polymer deposition for implantable applications (see later). Molecular templating can be used to advantage to facilitate incorporation of conducting polymers into textiles [41] (Figure 11.10). This approach introduces an additional step into the in situ process discussed above. The introduction of the molecular template, usually an anionic polyelectrolyte, facilitates the integration of some ICPs into textiles. The template plays a dual role in that it is initially integrated into the textile and then acts as a template for incorporation of monomer (e.g., aniline). After addition of oxidant to initiate polymerization, the template acts as a molecular dopant in the ICP. The presence of the template localizes the polymerization process within the textile and prevents it from occurring in the bulk solution. For example, polyaniline is not readily incorporated into wool structures, however, poly(2-methoxyaniline-5-sulfonic acid (PMAS) is (Figure 11.11). The PMAS can subsequently be used to template the aniline monomer before chemical oxidation. Biological molecules such as chondroitin sulfate, dextran sulfate, and DNA have been shown to be effective molecular templates for this purpose. Wearable biomedical devices may also be realized by integration of conducting polymer fibers. The production of stand-alone ICP fibers is at present limited to polyaniline [42]. The high concentration
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SO3−
SO3H +
H N
N H n
+
MeO
MeO
FIGURE 11.11 Structure of poly(2methoxyaniline 5-sulfonic acid).
solutions required for wet spinning can be prepared in solvents such as NMP or DMPU. More recently, dichloroacetic acid has been used to prepare the solvent feed. This has enabled the production of polyaniline fibers in the most conducting (emeraldine salt) form. Conductivities as high as 900 S cm1 with tensile strength of 40–60 MPa have been reported [43].
11.2.2 Review of Electronic and Redox Properties ICPs are also electroactive materials that readily undergo oxidation and reduction processes as simplistically described for PPy in Figure 11.12 for anion or cation incorporation or expulsion. For a given polymer composition, the ability to incorporate or expel anions or cations can be influenced by the mobility of the dopant anion, the chemical environment in which the polymer operates, or the nature of the electrical stimuli applied. A number of mechanical, electronic, and chemical transitions of relevance to biomedical applications accompany these redox processes. These include a change in the modulus of the materials [44], the electronic conductivity [45] and capacitance, which decrease upon electrochemical reduction. The surface energy, directly observable via contact angle measurements, also changes upon reduction. The nature and magnitude of these changes are critically dependent on the processes that predominate (see Figure 11.13). The oxidation or reduction process, if repeatedly cycled, with frequencies in the time domain of seconds also generates significant ion fluxes [46,47]. These dynamic properties can be used to great advantage in designing materials capable of monitoring and manipulating biological systems—from the molecular to the mechanical (skeletal) level.
11.2.3 Sterilization For these materials to be utilized in biomedical in vivo and in vitro applications, it is imperative that the material be sterilized before use. Sterilization of these materials can be performed using a variety of methods such as ethylene oxide gas [48–50], ultraviolet light [51,52], ethanol washing [53], and autoclaving [12,54]. An alternative approach to sterilization of the formed ICPs is the polymerization of the ICP under sterile conditions. This approach was utilized by Richardson et al. [55] to prepare PPycoated electrodes for use in auditory nerve studies.
+
0
A−
N H
n
+e −e
m
−
N H
+A n
m
For small mobile anions −A− + − A N H
n
A− X+
+e −e
m
N H
n
m
For large immobile anions −A , X+ is cation from electrolyte solution −
FIGURE 11.12
Electrochemical switching of PPy exhibiting anion or cation movement.
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Biomedical Applications of Inherently Conducting Polymers (ICPs)
0
+
A−
N H
n
+e m
Conductive Hydrophylic Higher modulus High ion exchange capacity
FIGURE 11.13
−e
−
N H
+A n
m
Less conductive Hydrophobic Low modulus Low ion exchange capacity
Some of the changes accompany oxidation=reduction of PPy.
Wang et al. [54] sterilized PPy using an autoclave technique and reported that PPy has a good biocompatibility with rat peripheral nerve tissue and that PPy might be a candidate material for bridging the peripheral nerve gap. Song et al. [52] used ultraviolet light as the sterilization method in the development of a new platform, based on PPy, for characterizing neurite extension in complex environments. Ethylene oxide gas was used by Zang et al. [50] to sterilize PPy-coated polyester. This study investigated the basic biocompatibility aspects of two types of PPy-coated polyester fabrics for possible use as vascular prostheses. Their findings suggested that cell adhesion moieties should be incorporated into the PPy or fabric composite to increase cell adhesion and subsequent cell proliferation. The studies presented above indicate that sterilization has no detrimental effect on the ability of ICPs to support the growth of a range of cell types. However, very little is known about the effects of sterilization on the electronic and redox properties of the ICPs. One paper by Campbell et al. [12] showed that sterilization using autoclaving had little effect on the ICP conductivity. The conductivity and redox properties of ICPs are two features that make them an attractive material in biomedical applications. For example, the inherent electrical conductivity of ICPs opens up the possibility of altering the cell-growth characteristics of ICP-based materials by the application of electrical stimulation. Therefore, it is crucial to evaluate their electrical and redox properties subsequent to sterilization.
11.3
Biomolecular Sensing
The use of conducting polymers as biosensors has been extensively reviewed [56–58]. Although practical applications of such devices are yet to be realized, research within the field highlights some important principles (and knowledge) that can be utilized in other areas. Each biosensing system requires the presence of a biomolecular recognition site and a means of signal transduction. It has been found that a wide range of biological entities from enzymes to antibodies to entire living cells can be incorporated either as dopant anions A or by direct covalent attachment. It has also been shown that changes in electronic properties can be transduced to the outside world with accompanying biomolecular events. In addition, it has been shown that at least some biomolecular events (even highly specific antibody– antigen [Ab–Ag] interactions) can be influenced by the application of electrical stimuli.
11.3.1 Oligonucleotides Bidan and coworkers utilized an elegant approach to biosensor development that involved covalent attachment of oligonucleotides to the NH group of a pyrrole monomer (Figure 11.14) [59]. From this, a novel electrospotting technique that enabled arrays of bioactive dots was developed [60]. Molecular binding events were monitored using a range of transduction methods including electrochemical, quartz crystal microbalance, surface plasma resonance, and fluorescence [61].
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NH CH3
N n (CH2)2
NH
NH(CH2)2N N
HO
O
NH
m
NH O
N
NH
e− H+
NH
CH3
+ electrochemical oxidation
O O P O−
HO
N O
N
O
Oligonucleotide
O
O O P O− O
Oligonucleotide
FIGURE 11.14 Electrosynthesis of PPy copolymers containing oligonucleotide. (From Livache, T., Roget, A., Dejean, E., Barthet, C., Bidan, G., and Teoule, R., Nucleic Acids Res., 22, 2915, 1994.)
Garnier et al. [62] have also utilized oligonucleotides as the molecular recognition site. The oligonucleotides were covalently bound to the pyrrole monomer. This was achieved by reacting 3-N-hydroxyphthalimide pyrrole with 3-carboxyethyl pyrrole to produce an electroactive copolymer attached to an electrode surface (Figure 11.15). The N-hydroxyphthalimide then functioned as an easily displaced leaving group to an amino-substituted oligonucleotide. This biosensing system demonstrated a remarkable capacity to differentiate between the length of base sequences in complimentary binding oligonucleotide. An alternative approach to covalent immobilization involves integration of the oligonucleotide as the dopant into an ICP during electropolymerization. This approach introduced by Wang and coworkers [63,64] utilized a novel inverted electrochemical cell that enabled electropolymerization to be carried out in extremely small (25 mL) volumes (Figure 11.16). In these previous studies, both constant potential amperometry [63] and potentiometric stripping analyses [64] were used to monitor hybridization. Using oligonucleotides, incorporated as the molecular dopant, it has been shown that pulsed ampero-
O O
OH
O
O
O
N O
O
OH
O
O
N O
electropoly 0.6
N H
+0.4
N H
N H
0.6
N H
0.4 n
O O
OH
O
ON
O
O
OH
O
H CCT−AAG−AGG−GAG−TG N
H2N-ODN N H
0.6
N H
0.4 n
N H
0.6
N H
0.4 n
FIGURE 11.15 An alternate route to PPy copolymers. (From Garnier, F., Korri-Youssoufi, H., Srivastava, P., Mandrand, B., and Delair, T., Synth. Met., 100, 89, 1999.)
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Reference electrode (3.0 M Ag/AgCl)
metric detection (using controlled pulsed potentials) can be used to control the biomolecular interactions of interest in a flow injection analysis system [65].
11.3.2 Enzymes A range of enzyme-based sensors has been developed over the past decade. Although elegant synthetic routes to covalent attachment have been developed (see above), the convenience of direct incorporation into the conducting polymer backbone at the time of synthesis has attracted most attention. For example, Adeloju and coworkers have used direct incorporation of enzymes during electrosynthWorking electrode esis to develop enzymatic biosensors for detection of urea [66], sulfite [67], glucose [68], and formate [69]. Evidence for the direct electron transfer from PPy to an entrapped quinohemoprotein alcohol dehydrogenase (QHADH) from Gluconobacter sp.33 prepared via an in situ FIGURE 11.16 Schematic diagram of the polymerization of pyrrole in the presence of QH-ADH has setup of a micro electropolymerization cell. been demonstrated by Ramanavicius et al. [70]. It was A drop of solution was placed on the glassy proposed that the cooperative action of the pyrroloquinocarbon electrode surface. line-quinone (PQQ) and heme-containing enzymes permit electron transfer from the enzyme active site to the ICP. Ethanol is said to diffuse to the PQQ enzyme centre where it is oxidized to an aldehyde. The PQQ centre is subsequently regenerated by the heme sites in the enzyme. Resulting is an electron that can be readily transferred to the PPy at a viable kinetic rate. In another approach, Zeng et al. [71] developed an amperometric detection system for histamine using a methylamine dehydrogenase (MADH) Ppy-based sensor, which was found to be suitable for the detection of primary amines. Electrocatalytic films were prepared by depositing a polymer coating from an electrolyte solution containing MADH, pyrrole, and potassium ferricyanide and then finally dip coating in a Nafion (registered trademark of E. I. du Pont de Nemours and Co., Wilmington, Delaware). The potassium ferricyanide was required to keep the MADH in the oxidized form and to act as an electron transfer mediator from the enzymatic centre to the electrode surface (Figure 11.17). In the absence of ferricyanide, the sensor was unresponsive to histamine and other primary amines. The Nafion membrane functions as a hydrophobic barrier such that the hydrophilic methylamine had no sensor response whereas the hydrophobic histamine, propylamine, and butylamine did. The sensor was nonresponsive to glutamine, creatine, urea, and ascorbic acid at concentrations which are normally Platinum wire auxiliary electrode
PPy + Nafion R−CH2NH2 H2O
TTQ (MADH)
2 Fe(CN)64− 2e−
R−CHO +NH3 +H+
TTQH (MADH)
2 Fe(CN)63−
Electrode
FIGURE 11.17 Electrocatalytic film sensor construction with a hydrophobic Nafion barrier to provide histamine sensing selectivity.
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Conjugated Polymers: Processing and Applications
found in blood samples. The authors suggested that such a sensor would be useful for detecting neoplastic marrow diseases, polycythemia, and chronic myelogenous leukemia where histamine levels in blood are elevated well above the normal 0.1–1 mM level. Asberg and Inganas [72,73] have investigated the use of hydrogel ICP composites to form 3-D enzyme electrode sensors. In these studies, a Baytron-P (registered trademark of H.C. Starck GmbH, Goslar), poly(ethylenedioxythiophene) (PEDOT) stabilized by polystyrene sulfonate (PSS), was mixed with poly(4-vinylpyridine) as a cast film, which was ionically cross-linked by MgSO4. To add biofunctionality, biomolecules were added to the solution before casting and cross-linking. Horseradish peroxidase (HRP), bovine serum albumin, glucose oxidase, or fibrinogen were readily incorporated using this approach. To regenerate the enzyme after a recognition event, osmium (K2OsCl6) was also incorporated into the film as an electron transfer mediator between the enzyme and the ICP matrix. Using osmium, detection of HRP was achieved at 0.05 V. This low detection-potential was reported to be beneficial for improving the signal-to-noise ratio as it helped minimize current responses from common redox-active interferences. A number of strategies that involve incorporation of HRP into polyaniline have been described. Bartlett and coworkers [74] achieved this either via simple adsorption of HRP onto electrodeposited polyaniline or by electrodeposition of an insulating poly(1,2-diamino benzene) film containing HRP onto polyaniline. Smyth and coworkers [75] have shown that immobilization of HRP into polyaniline can be achieved by electrostatic attraction. It has recently [76] been shown that the ability to nanostructure the polyaniline layer has a significant effect on the ability to subsequently immobilize the protein and then on sensor performance. Others [77] have immobilized HRP into a sulfonated polyaniline–polylysine complex. This approach is interesting in that all of the individual components are water soluble (poly(5-methoxyaniline-2-sulfonate (PMAS), polylysine, and HRP) and this could lead to some novel processing (fabrication routes).
11.3.3 Antibodies and Antigens As discussed above, strategies that enable covalent attachment of Ab recognition sites to conducting polymers have been developed. However, all evidence suggests that it is the intimate association brought about by incorporation of the Ab as the dopant molecule or even physical entrapment that enables effective electronic communication with the biomolecules. This in turn allows changes in electronic properties arising from Ab–Ag interactions to be detected and allows for the Ab–Ag interaction to be manipulated via electrical stimulation. In the early 1990s, a label-free electrochemical method for detecting Ab–Ag interactions was developed by immobilizing the Ab into the ICP at the point of polymer synthesis [78–80]. The immunosensor was fabricated by entrapping antibody into a PPy film using electrodeposition in the presence of the antibody. The antibody-modified electrode was used in a flow injection apparatus with a pulse electrochemical detector (PED), which continually and rapidly switched the conducting polymer between the reduced and oxidized forms. Injection of a sample containing the antigen resulted in a well-defined, repeatable electrochemical signal, which was dependent on the concentration of antigen in the sample. There are two very surprising aspects to this behavior. Firstly, an electrochemical signal was observed from the injection of a sample despite the only change at the sensor—the binding of a redoxinactive antigenic species. Secondly, the return of the sensor signal to baseline and repeatability of the signal over more than 10 injections of antigen suggested that the antigen–antibody binding was reversible. The reversibility of the binding was surprising in view of the high affinity of antibodies for their target antigen (affinity constants typically of the order of 108–1012 M1), which from a practical point of view makes the binding irreversible. The practical irreversibility of the antibody–antigen binding equilibria is one of the principles upon which immunoassays are based as otherwise the separation and washing would result in dissociation of the immunocomplexes. However, the reversible nature of the process developed using electrical stimulation of ICPs is extremely useful, from the point of
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Biomedical Applications of Inherently Conducting Polymers (ICPs)
view of a reusable sensing technology. Since the initial reports, this system has been shown to be useful for detection of a variety of proteins, such as human serum albumin [78,80], p-cresol-BSA conjugates [79], thaumatin [81], polychlorinated biphenyls [82], and isoproturon [83]. Application of the repetitive pulsed potential results in the polymer oscillating from the oxidized to the reduced form. Given the immobile nature of the Ab functioning as the molecular dopant, then cations from the supporting electrolyte will be continuously incorporated or expelled as the polymer oscillates between these oxidation states (Figure 11.18). This ion influx and efflux as well as the change in hydrophobicity or hydrophyllicity of the polymer backbone as it is oxidized or reduced obviously have a dramatic effect on the Ab–Ag binding reaction. The electronic signal arising is attributed to the fact that the inherent capacitance of the PPy-Ab system is altered when the Ag binds. This change in capacitance results in a change in charging current as the potential is pulsed giving rise to a direct electronic signal from the Ab–Ag binding event. Sargent and Sadik [84] have proposed that there are four steps involved in the current generation at an antibody-immobilized conducting polymer electrode: 1. 2. 3. 4.
Charge transfer at the Ppy–electrode interface Migration of ions through the polymer membrane to balance charge Diffusion of ions from the solution to the electrode Adsorption or desorption of the antigen at the PPy and solution interface
The slow rate of adsorption or desorption in Step 4 was considered by Sargent and Sadik [84] as the ratedetermining step, which is consistent with the antigen binding causing a change in the recorded current. Impedance spectroscopy measurements show that a change in potential from the reduced state to the oxidized state is accompanied by an increase in the capacitance of the electrical double layer, Cdl [85]. On exposure to the antigen, however, a smaller increase in Cdl was observed. There was also a change in charge transfer resistance (R CT) that Gibson et al. [86] believe is the dominant factor in the change in the total impedance observed with antigen binding. Thus, the pulsing of the polymer between its oxidized and reduced states results in fluxes of ions into and out of the polymer.
+
N
Ab−
+0.4V (− e−)
n
H
(a)
+
H
n Na+
(b)
Antigen binding promoted
N
Ab−
n
H
Cation influx
Cation efflux
N
0
−0.2V (+ e−)
Ab− Ag−
0
−0.2V (+ e−) +0.4V (− e−)
Ab− Ag−
N
n
H Na+ Antigen binding hindered
FIGURE 11.18 The proposed mechanism for signal generation in the PED immunosensor. All voltages are related to the Ag=AgCl reference electrode system.
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Conjugated Polymers: Processing and Applications
Another interesting example involving ICPs, Abs, and Ags involves the detection of Anti-D biomolecule important in analyses of blood sera. In this case, the D antigen needs to be incorporated within the ICP. The D antigen is distributed throughout red blood cell membranes; hence, intact red blood cells were incorporated into ICPs during electrodeposition [87,88]. This work also highlighted the importance of polyelectrolyte dopants acting to facilitate the red cell incorporation into the ICP and to provide a hydrogel like environment within the ICP matrix for the cells to continue functioning.
11.4
Biomolecular Actuators
ICP-based controlled release devices rely on the principle of release of a biologically significant compound from the polymer matrix under the influence of applied potential. Huang et al. [89] produced a PPy film electrode doped with 5-fluorouracil (5-FU), an anticancer drug. It was shown that incorporated 5-FU will slowly undergo ion exchange with Cl and then application of 0.6 V versus Ag=AgCl resulted in a rapid release of the anticancer drug. The rate of release was found to be dependent on the applied reduction potential and the polymer film thickness. The use of PPy membrane as an ion gate for the controlled release of anionic drugs based on salicylate, naproxen, and nicodise was demonstrated by Sundholm and coworkers [90]. Using 1–6 mm films, less than 5% of the bound anion was released from the membrane by ion exchange, however, on electrochemical reduction the anionic drugs were freely released. The amount of drug stored within the film was dependent on the molecular size, the drug structure, and the method used for polymer synthesis. Pernaut and Reynolds [91] have investigated the use of PPy as a model material, which has the ability of sensing and responding by releasing a biologically active molecule. In this study, adenosine triphosphate (ATP) was employed as a model drug, which was loaded into the polymer film during electrochemical synthesis. Both chemical (hydrazine reduction) and electrochemical triggers were employed with release rates for the latter of up to 20 mg cm2 min1 from a 10 mm thick conducting polymer membrane. Diffusion rates of the ATP from the polymer during stimulation at 0.7 V versus Ag=AgCl of 5 109 cm2 s1 were determined. Massoumi and Entezami [92] reported the controlled release of dexamethasone sodium phosphate (DMP) from a conducting polymer bilayer film consisting of a PPy inner film doped with DMP and poly(N-methylpyrrole)=polystyrene sulfonate (PNMP=PSS) or polyaniline sulfonate (SPANI) outer film. DMP was released from the inner film by an application of less than 0.6 V. In this device, the outer polymer layer functions as an ion and solvent barrier and also effectively reduces the rate of DMP release under an applied reducing electrochemical field, thereby providing an additional route to controlling release rates. It has been demonstrated that it is possible to perform controlled release of biologically active molecules utilizing the actuation of ICPs as a flow-gating device. Low et al. [93] have developed the concept of responsive controlled drug release utilizing the actuation capabilities of a ICP, namely polyaniline and a polyaniline composite with poly(2-hydroxyethylmethacrylate) hydrogel in combination with a solid-state silicon-based microvalve structure with metallic electrode contacts (Figure 11.19). A number of microvalve configurations have been proposed from a sphincter arrangement where the ICP element acts in a constricting fashion, a plunger to a constricting tube arrangement. It was demonstrated that these arrangements acted reversibly, whereas metallic barrier layer valves acted as one-shot devices. Polyaniline–hydrogel blends demonstrated superior actuation properties than the pure polyaniline under the microvalve test conditions. This was primarily due to the hydrogel possessing a higher degree of swelling, which combined with the redox switchability and associated volume changes in the conducting polymer gave clear advantages over the hydrogel and ICP alone. Potential applications for these device structures for use as a drug delivery system (reservoir) and biosensor under microprocessor control were discussed. Perhaps one of the most elegant potential applications of conducting polymer actuators in the biomedical field is the possibility of creating cellular microhospitals that could encapsulate cells and treat them via controlled release of drugs [94].
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Drug flow Insulator
Insulator
M
cle
us c
us
le
M
Conductor Insulator
Conductor Insulator
Body fluids (a)
Drug flow Insulator
Insulator
Insulator
Insulator Muscle
Access holes
Body fluids
Access holes
(b)
Insulator
Insulator Drug flow
Conductor Muscle Constraint tube
Conductor Muscle Access holes Body fluids (c)
FIGURE 11.19 Schematic of alternate designs for opening and closing holes in a drug reservoir using the artificial muscle concept. (a) Sphincter configuration (Design A); Plunger configuration (Design B); and (c) Tube configuration (Design C). (Designs proposed by Low, L-M., Seetharaman, S., He, K-Q., and Madou, M.J., Sens. Actuators B, 67, 149, 2000. With permission.)
11.5
ICPs in Tissue Engineering
The goal of accurately mimicking the structure and functional profile of the materials found in the body’s native extracellular matrix has been sought since the early days of surgeon Johann Friedrich Dieffenbach, who performed tissue transplantation experiments in 1822 [95]. Although cells have been cultured, or grown, outside the body for many years, the possibility of growing complex, threedimensional tissues—literally replicating the design and function of human tissue—is a recent development. Tissue engineering (TE) has been recognized for sometime as a promising alternative to donor
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tissues, which often are in short supply and are susceptible to immunological problems associated with infectious disease. The hope is that the biological function lost in the host tissues will be able to be restored and maintained by TE. Advances in TE now depend on developments in material science, nanotechnology, and bio-nanotechnology. The emergence of many future TE technologies will ultimately require biomaterials to be designed to support tissue growth physically, as well as to elicit desired receptor-specific responses from particular cell types [96]. To date nonelectroactive polymers such as polylactide [97], polyglycolide [98], polycaprolactone [99], and their copolymers have been used extensively in TE applications [100,101]. The prospect of using materials that incorporate stimulatory cues, such as electrical signals, to regulate cell attachment, proliferation, and differentiation [102] is an appealing one. Fabricating TE scaffolds that are electrically conducting would allow for the use of electrical stimulation to induce these processes and ultimately promote protein secretion to encourage the formation of a natural extracellular matrix. Stimulus-responsive polymers show great promise as new materials in biomedicine [103]. ICPs are stimulus-responsive polymers that can be synthesized to form composites that could serve as smart biomaterials. ICPs are ideal candidates for use in TE applications as they allow precise external control over the level and duration of stimulation. PPy was identified, as early as 1994, as a suitable candidate that may yet revolutionalize TE by facilitating the proliferation of mammalian cells [104,112]. The combination of PPy with appropriate polyelectrolytes results in a biomaterial with unique electronic properties and controlled release capabilities. Polypyrrole is perhaps the most widely studied ICP for TE applications because of its chemical and thermal stability, ease of preparation and electroactivity [105], and its in vitro compatibility [104]. Polypyrrole doped with nonbiologically active dopants, such as tosylate (pTS), have been characterized for biological interactions as they can trigger cellular responses in biological applications [106]. However, incorporation of more biologically active dopants has shown greater promise. Polypyrrole can be synthesized using a variety of negatively charged dopants that can be chosen for their biological significance such as HA and heparin [107,108]. Polypyrrole has been synthesized to form composites that could serve as smart biomaterials, such as animal tissue hybrids [109–111], substrates to support the growth of several mammalian cell types [104,112,113], and provide bioelectric fields in cultures of nerve cells [114]. Collier et al. [15] synthesized PPy in the presence of HA to form a biocompatible ICP that was evaluated for in vitro cell compatibility and tissue response in rats. The authors showed that the smooth PPy-HA films retained the HA in vitro for several days and promoted vascularization in vivo. The authors concluded that the PPy-HA composite biomaterials are promising candidates for TE and wound-healing applications that may benefit from both electrical stimulation and enhanced vascularization. Cen et al. [131] also demonstrated an in vitro bioactivity of HA-functionalized PPy with PC12 cell culture. Significant enhancement of PC12 cell attachment was noted in the presence of nerve growth factor, making this substrate a good candidate for nerve regeneration and repair applications. The ability of PPy to noninvasively control the shape and growth of mammalian cells was investigated by Wong et al. [104]. Their in vitro studies demonstrated that when the polymer was in the oxidized state, aortic endothelial cells spread normally on the surface and synthesized DNA. However, when the polymer was electrochemically switched to its neutral state both cell extension and DNA synthesis were inhibited without affecting the cell viability. Zhou et al. [14] and then Garner et al. [115,107] successfully polymerized PPy containing heparin as the dopant and evaluated its suitability for use as a substrate to support endothelial cell proliferation. When PPy was grown in the presence of heparin the resulting polymer can sustain the attachment and proliferation of human umbilical vein endothelial cells (HUVEC). The HUVECs spread and adopted morphology typical of that grown under standard in vitro conditions on tissue culture polystyrene. However, when the PPy was grown in the presence of nitrate ions (NO3) the seeded HUVECs did not attach. This result was at least in part attributed to the increased hydrophobic nature of the PPy-NO3 compared to the PPy-heparin polymer.
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Li et al. [108] investigated the use of conducting polymers with heparin covalently immobilized onto the surface. Covalently bound heparin on poly(ethylene glycol)methacrylate (PEGMA)-functionalized PPy surfaces showed an increased biocompatibility at PEGMA concentrations of 5 vol %. Immobilized heparin resulted in a suppression of platelet adhesion and activation as well as a prolonged plasma recalcification time. Electrochemical stimulation also showed a beneficial effect. In a further study, Li et al. investigated protein adsorption and thrombus formation on heparin or PEGMA-functionalized PPy surfaces under electrical stimulation [116]. PPy heparin PEGMA composites displayed the highest ratio of albumin to fibrinogen adsorption, the lowest plasma absorption, and lowest thrombus formation. Under electrical stimulation (50 mA constant current), these processes were significantly reduced although 30–60 min prestimulation was required to achieve this effect. Extended film electroactivity was observed over long-term immersion (4 d), suggesting potential uses in intravascular applications. The tissue response to PPy-coated polyester fibers was investigated by Alikacem et al. [117]. PPycoated and -uncoated polyester fibers were implanted subcutaneously in rats for several days with the results revealing a more persistent tissue reaction for the most conducting PPy-coated material and a shorter acute tissue response as the surface resistance increased. Blood monocyte activation studies of tissue response to the implanted PPy-coated polyester indicated that the thickness of the PPy coating, which correlates with the conductivity, was directly related to the tissue response. The best result for the PPy-coated polyester fiber (i.e., lowest tissue response) was similar to the tissue response results for the uncoated material. The authors identified the PPy-coated polyester fiber as a suitable candidate for further studies. The in vivo biocompatibility and biostability of PPy-coated polyester fabrics was further examined by Jiang et al. [118]. Three PPy-coated fabrics were prepared using phosphorylation (PPy-Phos), plasma activation (PPy-Plas), and plasma activation plus heparin treatment (PPy-Plas-HE). Untreated and fluoropassivated fabrics (F-PET) were controls. The specimens were implanted subcutaneously in the back of rats for 3–90 d, then harvested and processed for enzymatic, histological, and morphological analyses. A noninvasive MRI method was used to continuously monitor the inflammation. The level of acid and alkaline phosphatase showed a similar or a less intensive cellular reaction to the PPy-coated fabrics, when compared to the controls. Histology supported the enzymatic results and showed a fast collagen infiltration at 28 d for the PPy-Phos fabric. MRI reported an overall decrease of inflammation over time, with the PPy-coated fabrics showing a similar or mild inflammation in contrast to the noncoated fabrics. PPy clusters and excessive PPy laminary-coating on the PPy-Plas and PPy-Plas-HE were lost with the implantation. This experiment suggests a similar in vivo biocompatibility of the PPycoated and noncoated polyester fabrics and the importance of achieving a thin adherent uniform PPy coating. Polypyrrole–polyester (Milliken Style 205 doped with 1,5-naphthalenedisolfonic acid) fabric cellular responses were also investigated by Jakubiec et al. [119], who indicated that high levels of electrical conductivity (100–200 V=square) was detrimental and resulted in low cell migration, proliferation, and endothelial cell viability. The coated textile system was shown to influence both the morphology and the function of mammalian cells in vitro. Surface functionalization to improve polymer substratum biocompatibility has been examined by several research groups [120,121]. A surface modification technique was developed by Li et al. [122] for the covalent immobilization of heparin onto PPy films. Physicochemical and blood compatibility characterization of these PPy-modified films showed an improvement of film wettability while retaining significant electrical conductivity. With heparin immobilized, platelet adhesion and platelet activation was significantly suppressed, and the plasma recalcification time (PRT) was significantly prolonged. The authors showed that electrical stimulation plays a positive role in decreasing platelet adhesion and increasing PRT on pristine and surface-modified PPy films. These results are desirable to ensure the nonrejection of an implanted device. Lian et al. [123] investigated the polysaccharides, hyaluronic acid (HA) and sulfonated hyaluronic acid (SHA), as a suitable biomolecule to improve the surface biocompatibility of PPy films. The biological activity of the HA-functionalized PPY film was assessed by means of an in vitro PC12 cell
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culture. Cell attachment on the HA-functionalized PPy film surface was significantly enhanced in the presence of nerve growth factor. The PRT observed from the SHA-functionalized PPy film was significantly prolonged compared with the HA-functionalized PPy film. Some reduction of platelet adhesion was observed for the SHA-functionalized PPy film, compared with that of the HA-functionalized PPy film.
11.6
ICPs Used in Nerve Cell Regeneration
Electrical stimulation has been shown to enhance nerve cell regeneration [124,125], the mechanisms for this effect are, however, unclear. One hypothesis is that an electrical stimulus alters the local electrical fields of extracellular matrix molecules, changing protein adsorption [126]. As early as 1994, studies into the suitability of ICPs such as PPy as neuronal scaffolds provided positive proof that these electroactive stimulus response polymers indeed have a role to play. Shastri et al. [127] showed that neurite extension of PC 12 cells was more pronounced on PPy surfaces as compared to tissue culture polystyrene. The authors also showed that the application of an electrical stimulus to the cell culture on the PPy film significantly increased the expression of neurites in the cells compared to the controls. In addition, they demonstrated through tissue compatibility and transected sciatic nerve regeneration studies in rat models that the PPy films invoke little negative response and support nerve regeneration. Schmidt et al. [128] proposed that PPy could prove potentially useful not only in providing neuronal guidance but also in localizing electromagnetic stimulation when used as a polymer scaffold or guidance channel. The initial stages of their research showed that PC12 cells and primary chicken sciatic nerve explants, attached and extended neurites equally well on both PPy films and tissue culture polystyrene in the absence of electrical stimulation. However, PC12 cells cultured on PPy films showed a significant increase in neurite lengths compared with those not subjected to electrical stimulation. Kotwal et al. [126] used the electrically conducting polymer PPy to investigate the hypothesis that electrical stimulation increases the adsorption of serum proteins, specifically fibronectin (FN), thereby increasing neurite extension. The authors looked at the effects of electrical stimulation on protein adsorption when an electrical current was applied to PPy (a) during protein adsorption (immediate stimulation) and (b) several hours after protein adsorption (delayed stimulation). They found that immediate stimulation of PPy increases FN adsorption from purified FN and serum-containing solutions. Correspondingly, PC12 cells grown on PPy films that had been previously adsorbed with FN during immediate stimulation expressed longer neurites. However, for delayed stimulation, no significant differences in adsorption or neurite outgrowth were observed. These studies shed some light on the intimate interaction occurring at the polymer–neurite interface and suggest that increased FN adsorption with immediate electrical stimulation may explain enhanced neurite extension on electrically stimulated PPy. The effect of synthesis conditions on the suitability of a PPy scaffold to enhance neurite growth was investigated by Cui et al. [129]. Cui and workers precisely deposited PPy combined with biomolecules, having cell adhesion functionality, electrochemically onto microelectrode sites of neural probes. Their findings suggest that the nature of the grown PPy film has a bearing on the extent of neurite attachment and growth. A convex fuzzy morphology of a coating seemed to be better than a thinner and flatter surface for establishing an intimate interfacial contact with tissue. They also showed that a coated electrode had a much higher surface area and charge capacity, which provided a larger more efficient interface for electronic and ionic signal transport. PPy doped with conventional dopants such as sodium dodecylbenzenesulfonate [48] and sodium poly(styrenesulfonate) [128] have proved successful in neurite studies. The use of biologically significant dopants has also shown considerable improvements in neurite attachment and extension. When a protein polymer containing fibronectin fragments (SLPF) and a nonapeptide (CDPGY1GSR) were used as the PPy dopant, the cellular response was improved compared to CH3COO doped
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PPy [129]. Glial cells appeared to attach better to Ppy- and SLPF-coated electrodes than uncoated electrodes, whereas neuroblastoma cells grew preferentially on and around the Ppy- and CDPGY1GSRcoated sites. The PPy-CH3COO coating on the same probe did not show a preferential attraction to the cells. PPy has dominated the research activities concerning the use of the ICP in neuronal scaffolds, but investigations into the feasibility of other ICPs, such as polyaniline (PANI) has also been performed. Oren et al. [130] reported on a novel route to synthesize 2D-polyaniline (2D-PANI) on sulfonatedpoly(styrene) (SPS) templates. Culture studies showed that Aplysia neurons grown on 2D-PANI exhibit an unusual growth pattern and adhesion to this conducting substrate that is manifested by the formation of giant lamellipodia. This behavior is characteristic to uniform substrates containing only 2D-PANI. However, in patterned substrates containing additional poly(L-lysine) Aplysia neurons prefer to extend new neurites on the poly(L-lysine) domains. Polypyrrole implants for neural prosthetics, namely brain tissue, have been investigated by George et al. in 2005 [48]. In this study, PPy doped with polystyrene sulfonate or dodecylbenzenesulfonate were implanted into a rat cerebral cortex and shown to perform as well as or better than Teflon with neurons and glial cells enveloping the implant providing a direct electrical route to brain parenchyma via the implants. It is proposed that such implants may provide a route for the transmission of external and internal electronic signals to brain tissue postoperatively to similar and repair damaged neural structures. It has been shown previously that nerve axions can grow on a PPy surface and that the development of axions can be influenced by controlled release of nerve growth factors doped within the polymer to stimulate axion formation [131]. In a study by Cui et al. [129], it was also shown that Ppy-coated electrodes could effectively transduce AC-modulated electrical stimulus to neural tissues via in vivo recording. The biocompatibility of PPy therefore makes it an exciting substrate for neural scaffold structures, nerve stimulation by charge injection, and implant devices. A proposed cochlear implant device involves the combination of piezoelectric PVDF films with PPy coatings [132]. In this case, the PVDF acts to transduce sound pressure waves into an electrical signal at the external electrode contacts, which then transmit the generated electrical charge directly to the nerve fibers within the cochlear, via conducting electrodes on the PVDF surface. Traditional metallic electrode contacts at the surface of the piezoelectric PVDF typically have an acoustic impedance mismatch with the biological, fluid filled environment found within the cochlear. In situ deposition of thin film PPy coatings onto the PVDF sensor provided an improvement in device performance as it was flexible, provided a low acoustic impedance and elastic modulus match to the biological environment while simultaneously functioning as a conductor for charge transmission to the nerve cell. Incorporation of ICPs into hydrogel has been considered a suitable route for further compatibilizing the polymer as the hydrogel itself is a well-known and accepted biomaterial in pharmaceutical and biomedical applications, such as tissue scaffolds, drug delivery agents, etc., and once modified has been reported to have similar mechanical properties as brain tissue. It is this capability that lead to an investigation by Kim et al. [133], who looked at PPy hydrogel scaffolds coated onto microfabricated neural prosthetics devices. It was shown that the PPy hydrogel scaffold provided a lower electrical impedance (7 kV) than for PPy films alone (100 kV) at a biologically significant stimulation frequency of 1 kHz. This composite effectively provided a mechanical buffer between the soft brain tissue and the hard neural probe while also having a lower electronic impedance.
11.7
Biodegradability and Stability
An important aspect of implantable ICPs is biodegradability. Ideally, these materials will provide a suitable scaffold to promote cell or neurite attachment and proliferation; provide an electrical stimulation to a point where the cells or the neurites are self sustaining and then degrade away. The degradable products of the polymers need to be expelled from the body, leaving behind the body’s own tissue. Many researchers have been synthesizing PPy, which has the ability to degrade over time.
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Shi et al. [134] recently developed a novel biodegradable composite conducting polymer incorporating PPy and polylactide. This conducting biodegradable composite was investigated for its suitability as a substrate for electrically stimulated cell culture and as a scaffold to explore the potential of electrically stimulated tissue regeneration. A composite membrane containing 5% PPy (PPy=PLGA) showed promising results in an in vivo study. Wang et al. studied the in vivo biocompatibility performance of PPy=poly(D,L-lactide) composite and PPy-coated poly(D,L-lactide-co-glycolide) membranes implanted subcutaneously in rats for 3–130 d [49]. Inflammation responses were observed in the early implantation stages (day 3) but this response decreased with extended implantation time. Both implanted materials were observed to have no abnormal tissue responses with respect to commonly implanted control specimens. The biodegradable materials synthesized by Shi et al. [134] and Wang et al. [49] were a composite combining a conducting polymer with traditional biodegradable tissue-engineered polymers. An alternative approach to producing biodegradable ICPs has been put forward by Schmidt et al. [135] and Rivers et al. [102]. The Schmidt et al. [135] synthetic strategy consisted of tethering conductive pyrrole– thiophene oligomers together with degradable ester linkages using an aliphatic linker. Rivers et al. expanded on this synthetic process by removing the thiophene oligomer and replacing it with another pyrrole oligomer. The synthesis of these polymers has been successful, however to date no cell culturing studies have been performed. Another important issue is the stability of the ICP in a biological in vivo or in vitro environment, which itself can be exceptionally aggressive. There is a fine balance required between compatibility with the environment and the time of exposure or service life for any application or device. ICPs in the conductive form are generally in a partially oxidized state and therefore have the potential to either be oxidized or reduced further. This capability opens up the possibility for the ICP to function as a free radical scavenger in biological systems, as noted by Gizgavic-Nikolaidis et al. [136,137]. In these studies, the radical scavenging ability of soluble polyaniline grafted to lignin, poly(anilinesulfonic acid), and PPy were investigated against a,a-diphenyl-b-picrylhydrazyl (DPPH), a stable and persistent free radical, in order to establish the reducing strength of these polymers. The studied ICPs exhibited a clear capacity to quench the otherwise stable DPPH radical. The authors note that these polymers may assist in the reduction of radical species that can contribute to degradation of biological materials, such as lipid structures and DNA, and may prove to be useful in biological media or as implantable structures. Given that PPy as a class of materials has significant biocompatibility and functionality, Thomas et al. [138] have developed a range of water-soluble poly(3,4-alkenedioxypyrroles) (PXDOPs) as a potentially stable biomaterial. Oxidized PPy is unstable to reduction by weak biologically relevant reducing agents such as dithiothreitol (DDT) and glutathione. In a case study, oxidized poly(ethylenedioxypyrrole) was observed to have a reduction potential lower than 0.5 V versus Standard Calomel Electrode (SCE). These exceptionally low reduction potentials make the polymer resistant to strong biological reducing agents, thereby making them useful as a material in contact with biologically active environment, which requires the presence of oxidized, conductive polymer.
11.8
Biomechanical Sensing
An area that has been exploited for sometime is the application of ICP coated textiles in monitoring body motion. A review by Engin et al. [139] highlights some of the emergent developments and trends in biomedical sensors with a specific emphasis on intelligent textiles and wearable electronic devices. Devices such as Georgia Tech’s Wearable Motherboard (registered trademark of Georgia Institute of Technology, Atlanta, Georgia) and integrated switches such as the Softswitch (registered trademark of Softswitch Ltd, West Yorkshire, UK) are examples of how function is being embedded into everyday apparel. Embedding of wireless technologies into clothing with non-ICP-based textile sensors textrodes for biomedical monitoring, such as ECG and respiration rate functions in hospital environments, has been described by Catrysse et al. [140]. The extension into the ICP realm where there
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Resistance change (Rf−R0/R0*100%)
5 Strain %
0 −5
0
10
20
30
40
50
60
70
80
−10 −15 −20 −25 −30 −35 −40 −45 2Hz
3Hz
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FIGURE 11.20 Electrical resistance versus strain response over 1, 2, and 3 Hz frequencies of Nylon Lycra, coated with PPy-NDSA.
is a clear advantage in replacing metallic elements with polymeric elements has been realized by a number of groups. ICP-coated textiles prove to be highly effective wearable strain gauges. Most importantly, they have been shown to function well over the frequency ranges (1–3 Hz) and the strain ranges (10%–100%) needed for monitoring movement of joints. Typical strain responses over a range of frequencies are shown in Figure 11.20. The response is dependent on the base textile used with this providing a means to tune the sensitivity and response time.
11.8.1 Biomechanical Actuators: Artificial Muscles and Microactuators The use of conjugated polymer actuators in biomedical applications has been reviewed recently [141]. It is well known that the ICPs when oxidized or reduced undergo a dimensional change during redox cycling as a result of the incorporation or expulsion of anions or cations [142]. This process is a direct consequence of p-type doped ICPs having a positive charge on the polymer backbone in the oxidized (conducting) form, which generally results in the incorporation of an anionic species (dopant) to achieve electrostatic neutrality. If sufficiently large or sterically bulky anions are incorporated into the polymer during growth, this mechanism can be switched from predominantly anion incorporation or expulsion during redox cycling to that of cation incorporation or expulsion. Under these conditions, when the polymer is reduced, the polymer-backbone loses its cationic character but the anion cannot be expelled due to its being locked within the polymer matrix. The consequence of this is the preferential uptake of cations in the electrolyte solution. Upon reoxidation of the polymer, the cationic charge is reinstated upon the backbone and the unbound cation is subsequently ejected. By selective design an actuator device can be constructed which takes advantage of the associated volume changes within the polymer during redox cycling. The most common of these devices has been the bilayer actuator design wherein one layer in the device contracts or expands upon oxidation–reduction, resulting in a bending or curling type motion (Figure 11.21). A major limitation of such devices is that the materials used are not highly conductive and so there is a need for efficient electrical interconnects, if good efficiencies are to be realized. This problem is exacerbated by the fact that upon reduction the polymer conductivity is further diminished and this limits the electrochemical efficiency particularly in larger (greater than micron dimensions) devices
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Conjugated Polymers: Processing and Applications
TFSi− TFSi−
TFSi−
TFSi−
u+
TFSi− TFSi− u+
TFSi−
TFSi− TFSi− TFSi−
TFSi− TFSi−
TFSi−
u+ TFSi− TFSi− TFSi− TFSi−
u+
TFSi− TFSi−
TFSi− u+
u+ TFSi−
TFSi−
FIGURE 11.21 Schematic showing the bilayer actuator design, wherein one layer in the device contracts or expands upon oxidation–reduction resulting in a bending or curling type motion.
[143]. One approach to this problem is to use helical wire windings that function as a more efficient charge injection route to polymer actuators (Figure 11.22) [144]. The use of this helical interconnect results in improvements in electrochemical efficiency (percentage of polymer actually oxidized or reduced). It has recently been shown that incorporation of carbon nanotubes into polyaniline fibers [145] results in materials with improved actuation properties. Mazzoldi and De Rossi et al. proposed the fabrication of a steerable microcatheter device based upon a polyaniline fiber actuator encased in a solid polymer electrolyte (polyacrilonitrile in propylene carbonate and ethylene carbonate with a cupric chloride supporting electrolyte) [146,147]. The polyaniline fiber was prepared by coagulation spinning of emeraldine base, drawn and post-doped with HClO4. The resulting fiber is then encapsulated in the solid polymer electrolyte and a Cu wire coil counter electrode is wrapped around the inner structure. A steerable microcatheter can be fabricated by bundling these fiber actuators into the tip of a 1.65 mm surgical catheter with each fiber individually addressed by an activating potential to give direction control to the bending moment.
11.8.2 Wearable Prosthetics The design of solid-state actuator systems that can be incorporated into wearable fabric structures is one of growing interests both as a rehabilitation tool and biomechanical augmentation device. The use of ICP fibers that can be incorporated into textile structures as a sensor or actuating device for prosthetic applications has been reviewed previously [148]. Of particular interest is the development of a rehabilitation glove (Figure 11.23). The glove [149] is to be used to provide continuous passive motion during rehabilitation via physiotherapy after major injury and surgery to the hand. Applications such as these provide an excellent framework for development of the performance criteria of artificial muscle fibers. Fibers with length of the order of 30 cm with 5% strain capabilities under significant (5 MPa) load are required. The advent of helical wire interconnects,
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(1)
20
15
10
1 mm
0
(2) Epoxy Connector length
60 mm
Connector length A
B
C
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E
F
FIGURE 11.22 (1) Picture showing the hollow polymer actuator with helical interconnect at various diameters and (2) the schematic diagrams showing its construction method: (A) 25 mm of platinum wire is wrapped around the 125 mm wire as a spiral; (B) polymer synthesis—the assembly is placed in polymer electrolyte solution (0.5 M Py, 0.25 M TBA PF6 in PC) and electroplated for 24 h at 288C; (B) polymer coating forms around wire and spiral; (C) 125 mm centre wire is withdrawn from the polymer tube or helix; (D) two short connectors of 125 mm wire are inserted into each end; (E) 25 mm wire is pulled tight around these ends for a good electrical connection and epoxy glued to hold in place. (From Ding, J., Liu, L., Spinks, G.M., Zhou, D., Gillespie, J., Wallace, G.G., Synth. Met., 138, 391, 2003. With permission.)
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Conjugated Polymers: Processing and Applications
Sensor Patches
Actuating Fibers
FIGURE 11.23
Rehabilitation glove showing the positioning of the actuating fibers and sensing patches.
as discussed previously, as well as the use of ionic liquid electrolytes [150,151] have taken us some way toward these requirements with application in biomechanical monitoring and rehabilitation, and facilitating new training routines [152]. Utilizing ICPs in wearable apparel for electromagnetic irradiation shielding, antistatics, heating and cooling augmentation, and biomechanical feedback devices is a developing area of interest [153–155]. These applications rely purely on a resistive conducting polymer coating on a textile substrate. When the textile substrate is deformed, the coated yarn elements are either compressed or separated, resulting in a resistance drop and increase, respectively. This resistance change can then be monitored and calibrated to measure human motion. The challenge to date has been the development of air- and moisture-stable polymer coatings. Polypyrrole-coated Nylon-Spandex stretchable fabrics provide a unique method for making external contacts for electrotherapy to the human body [156,157]. As the textile can be tailored to fit the human form, there is a clear benefit from using the garment as a direct sensing or activating element. These PPycoated textile samples exhibited an improved electrical conductivity at up to 60% strain. Direct body contact at 40% strain exhibited good conformability, flexibility, and conductivity over 30 cycles with minimal IR heating during charging cycles of the electropad. Typical electrotherapy DC charges to the body are in the range of 4.5–45 V at 0.1–10 mA for nerve of muscle stimulation with an induction current of 2–10 mA for medical applications. Tests at up to 20 mA showed no degradation over 600 s. Eventhough the ICP-coated textile had a lower conductivity than a typical copper plated electrotextile transcutaneous electrical nerve stimulation (TENS) using the ICP material exhibited a significant TENS effect [157].
11.9
Conclusions and Future Developments: Bionics and Beyond
The application of conducting polymers at the interfaces between biology and electronics is an area of great importance. The need to extract information from the biomolecular scale through to the entire body, in real time, is an important and evolving science. Similarly, communication in the opposite
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direction to impart changes to the biological system of interest is also critical. Simultaneous implementation of both of these processes places us at the forefront of truly intelligent systems. At the simplest level, communication with amino acids, proteins, enzymes, antibodies, DNA, and whole cells is clearly important [158]. It is well known that biological systems can be influenced by the application of electric fields. The ICP provides us with the capability of charge injection or removal at biological interfaces due to its electronic conductivity. More interestingly, though is the redox capability of these polymers, whereby they can function as simple on–off switches, facilitate controlled release of molecular species of interest, or direct mechanical interaction with the biological interface in the form of actuation. As we approach the nanodomain, the properties of ICPs are dramatically altered [159]. It is also known that the introduction of nanostructure to material surfaces has a significant impact on protein and cellular interactions [160–162]. In one study involving the demixing of polystyrene and poly(4bromostyrene) to produce nanometre high islands, the nanotopography had the favorable effect of enhancing the spreading of endothelial cells across the nanotextured surface [163]. In another work, using nanophase ceramics an optimal grain size in the nanodomain was found to promote osteoblast adhesion [164]. Little work has been done on nano-biosurfaces based on ICPs, however, it is an area attracting great interest. For example, recent work has shown that proteins can be more effectively mobilized on a nanostructured polyaniline platform with a concomitant increase in biosensor performance [165]. It is known that more highly conducting nanodomains exist within the essentially amorphous ICP host structure and that electrochemical switching speeds can be extremely rapid in nanowires composed of these materials. It is also known that control of nanotopography in other biomaterials can have a profound effect on the adhesion of mammalian cells. Given the explosion of activity in the area of ICP nanostructures, there is no doubt that these enhanced properties should translate into more effective biomolecular sensors and actuators.
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116. Li, Y., K.G. Neoh, and E.-T. Kang. 2004. Plasma protein adsorption and thrombus formation on surface functionalized polypyrrole with and without electrical stimulation. J Colloid Interface Sci 275:488. 117. Alikacem, N., Y. Marois, Z. Zhang, B. Jakubiec, R. Roy, M.W. King, and R. Guidoin. 1999. Tissue reactions to polypyrrole-coated polyesters: A magnetic resonance relaxometry study. Artif Organs 23:910. 118. Jiang, X., Y. Marois, A. Traore, D. Tessier, L.H. Dao, R. Guidoin, and Z. Zhang. 2002. Tissue reaction to polypyrrole-coated polyester fabrics: An in vivo study in rats. Tissue Eng 8:635. 119. Jakubiec, B., Y. Marois, Z. Zhang, R. Roy, M.-F. Sigot-Luizard, F.J. Dugre, M.W. King, L. Dao, G. Laroche, and R. Guidoin. 1998. In vitro cellular response to polypyrrole-coated woven polyester fabrics: Potential benefits of electrical conductivity. J Biomed Mater 41:519. 120. Bousalem, S., C. Mangeney, M. Chehimi, T. Basinska, B. Miksa, and S. Slomkowski. 2004. Synthesis, characterization, and potential biomedical applications of N-succinimidyl ester functionalized, polypyrrole-coated polystyrene latex particles. Colloid Polym Sci 282:1301. 121. Altankov, G., V. Thom, T. Groth, K. Jankova, G. Jonsson, and M. Ulbricht. 2000. Modulating the biocompatibility of polymer surfaces with poly(ethylene glycol): Effect of fibronectin. J Biomed Mater Res 52:219. 122. Li, Y.L, K.G. Neoh, L. Cen, and E.T. Kang. 2003. Physicochemical and blood compatibility characterization of polypyrrole surface functionalized with heparin. Biotech Bioeng 84:305. 123. Lia, C., K.G. Neoh, Y.L. Li, and E.T. Kang. 2004. Assessment of in vitro bioactivity of hyaluronic acid and sulfated hyaluronic acid functionalized electroactive polymer. Biomacromolecules 5:2238. 124. Aebischer, P., R.F. Valentini, P. Dario, C. Domenici, and P.M. Galletti. 1987. Brain Res 436:165. 125. Valentini, R.F., T.G. Vargo, J.A. Gardella, and P. Aebischer. 1992. Biomaterials 13:183. 126. Kotwal, A., and C.E. Schmidt. 2001. Electrical stimulation alters protein adsorption and nerve cell interactions with electrically conducting biomaterials. Biomaterials 22:1055. 127. Shastri, V.R., C.E. Schmidt, H.T. Kim, J.P. Vacanti, and R. Langer. 1996. Polypyrrole—A potential candidate for stimuluted nerve regeneration. Mat Res Soc Symp Proc 414:113. 128. Schmidt, C.E., V.R. Shastri, J.P. Vacanti, and R. Langer. 1997. Stimulation of neurite outgrowth using an electrically conducting polymer. Proc Natl Acad Sci USA 94:8948. 129. Cui. X., V.A. Lee, Y. Raphael, J.A. Wiler, J.F. Hetke, D.J. Anderson, and D.C. Martin. 2001. Surface modification of neural recording electrodes with conducting polymer=biomolecule blends. J Biomed Mater Res 56:261. 130. Oren, R, R. Sfez, N. Korbakov, K. Shabtai, A. Cohen, H. Erez, A. Dormann, H. Cohen, J. Shappir, M.E. Spira, S. Yitzchaik. 2004. Electrically conductive 2D-PAN-containing surfaces as a culturing substrate for neurons. J Biomater Sci—Polym Ed 15:1355. 131. Cen, L., K.G. Neoh, Y. Li, and E.T. Kang. 2004. Assessment of in vitro bioactivity of hyaluronic acid and sulfonated hyaluronic acid functionalized electroactive polymer. Biomacromolecules 5:2238. 132. Dwivedi, A., and R. Roseman. 2003. In-situ development and study of conducting polymer electrodes on PVDF substrates for electro-acoustic application of cochlear implants. Mat Res Soc Symp Proc 771:123. 133. Kim, D-H., M. Abidian, and D.C. Martin. 2004. Conducting polymers grown in hydrogel scaffolds coated on neural prosthetic devices. J Biomed Mater Res 71A:577. 134. Shi, G., M. Rouabhia, Z. Wang, L.H. Dao, and Z. Zhang. 2004. A novel electrically conductive and biodegradable composite made of polypyrrole nanoparticles and polylactides. Biomaterials 25:2477. 135. Schmidt, C.E., T.J. Rivers, T.W. Hudson, and J. Collier. 2003. Modification of electroactive biomaterials for neural engineering applications. In: eds. J.F. Rubinson and H.B. Mark Jr., Conducting polymers and polymer electrolytes: From biology to photovoltaics. ACS Symposium Series 832, pp. 154–165. 136. Gizdavic-Nikolaidis, M., J. Travas-Sejdic, G.A. Bowmaker, R.P. Cooney, and P.A. Kilmartin. 2004. Conducting polymers as free radical scavengers. Synth Met 140:225.
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137. Gizdavic-Nikolaidis, M., J. Travas-Sejdic, G.A. Bowmaker, R.P. Cooney, C. Thompson, and P.A. Kilmartin. 2004. The antioxidant activity of conducting polymers in biomedical applications. Curr Appl Phys 4:347. 138. Thomas, C.A., K. Zong, P. Schottland, and J.R. Reynolds. 2000. Poly(3,4-alkyenedioxypyrrole)s as highly stable aqueous-compatible conducting polymers with biomedical implications. Adv Mater 12:222. 139. Engin, M., A. Demirel, E.Z. Engin, and M. Fedakar. 2005. Recent developments and trends in biomedical sensors. Measurement 37:173. 140. Catrysse, M., R. Puers, C. Hertleer, L. Van Langenhove, H. van Egmond, and D. Matthys. 2004. Towards the integration of textile sensors in a wire less monitoring suit. Sens Actuators A 114:302. 141. Smela, E. 2003. Conjugated polymer actuators for biomedical applications. Adv Mater 15:481. 142. Gandhi, M., P. Murray, G.M. Spinks, and G.G. Wallace. 1995. Mechanism of electromechanical actuation in polypyrrole. Synth Met 73:247. 143. Otero, T.F., and J.M. Sansinena. 1997. Bilayer dimensions and movement in artificial muscles. Bioelectrochem Bioenerg 42:117. 144. Ding, J., L. Liu, G.M. Spinks, D. Zhou, J. Gillespie, and G.G. Wallace. 2003. High performance conducting polymer actuators utilising a tubular geometry and helical wire interconnects. Synth Met 138:391. 145. Mottaghitalab, V., G.M. Spinks, and G.G. Wallace. 2005. The influence of carbon nanotubes on mechanical and electrical properties of polyaniline fibers. Synth Met 152:77. 146. Mazzoldi, A., C. Degl’Innocenti, M. Michelucci, and D. De Rossi. 1998. Actuative properties of polyaniline fibers under electrochemical stimulation. Mater Sci Eng C 6:65. 147. Mazzoldi, A., and D. De Rossi. 2000. Active polyaniline fibers for a steerable catheter. Proc SPIEEAPAD 3987:273. 148. Spinks, G.M., G.G. Walace, L. Liu, and D. Zhou. 2003. Conducting polymers electrochemical actuators and strain sensors. Macromol Symp 192:161. 149. The International Patent Application No. PCT=AU03=01138 filed 4 September 2003 entitled Movement facilitation device, inventors Timothy, R.D. Scott, Veronica A Vare, Peter Abolfathi, Gordon Wallace, Geoff Spinks, and Dezhi Zhou. 150. Lu, W., A.G. Fadeev, B. Qi, E. Smela, B.R. Mattes, J. Ding, G.M. Spinks, J. Mazurkiewicz, D. Zhou, D.R. MacFarlane, S.A. Forsyth, M. Forsyth, and G.G. Wallace. 2002. Use of ionic liquids for p-conjugated polymer electrochemical devices. Science 297:983. 151. Ding, J., D. Zhou, G. Spinks, S. Forsyth, M. Forsyth, D. MacFarlane, and G.G. Wallace. 2003. Chem Mater 15:2392. 152. Holcombe, B., and G.G. Wallace. 2002. The brave new world of wearable intelligence. Wool Tech Sheep Breed 50:312. 153. Dall’ Acqua, L., C. Tonin, R. Peila, F. Ferrero, and M. Catellani. 2004. Preformances and properties of intrinsic conductive cellulose-polypyrrole textiles. Synth Met 146:213. 154. Tessier, D., L.H. Dao, Z. Zhang, M.W. King, and R. Guidoin. 2000. Polymerization and surface analysis of electrically conductive polypyrrole on surface-activated polyester fabrics for biomedical applications. J Biomat Sci—Polym Ed 11:87. 155. Bhat, N.V., D.T. Seshadri, and S. Radhakrishnan. 2004. Preparation, characterization, and performance of conductive fabrics: Cotton þ PANi. Textile Res J 74:155. 156. Oh, K.W., H.J. Park, and S.H. Kim. 2003. Stretchable conductive fabric for electrotherapy. J Appl Polym Sci 88:1225. 157. Kim, S.H., K.W. Oh, and J.H. Bahk. 2004. Electrochemically synthesized polypyrrole and Cu-plated nylon=Spandex for electrotherapeutic pad electrode. J Appl Poly Sci 91:4064. 158. Kane-Maguire, L.A.P., and G.G. Wallace. 2001. Communicating with the building blocks of life using organic electronic conductors. Synth Met 119:39. 159. Innis, P.C., and G.G. Wallace. 2002. Inherently conducting polymer nanostructures. J Nanosci Nanotechnol 2:441.
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160. Valsesia, A., P. Colpo, M.M. Silvan, T. Meziani, G. Ceccone, F. Rossi. 2004. Fabrication of nanostructured polymeric surfaces for biosensing devices. Nano Lett 4:1047. 161. Boeckl, M.S., T. Baas, A. Fujita, K.-O. Hwang, A.L. Bramblett, B.D. Ratner, J.W. Rogers, and T. Sasaki. 1998. Template-assisted nanopatterning of solid surfaces. Biopolymers 47:185. 162. Curtis, A., and C. Wilkinson. 2001. Nanotechniques and approaches in biotechnology. Trends Biotechnol 19:97. 163. Dalby, M.J., M.O. Riehle, H. Johnstone, S. Affrossman, A.S.G. Curtis. 2002. In vitro reaction of endothelia cells to polymer demixed nanotopography. Biomaterials 23:2945. 164. Webster, T.J., R.W. Siegel, and R. Bizios. 1999. Osteoblast adhesion on nanophase ceramics. Biomaterials 20:1221. 165. Morrin, A., O. Ngamna, A.J. Killard, S.E. Moulton, M.R. Smyth, and G.G. Wallace. 2005. An amperometric enzyme Bbosensor fabricated from polyaniline nanoparticles. Electroanalysis 17:423.
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12 Biosensors Based on Conducting Electroactive Polymers 12.1 12.2
Introduction..................................................................... 12-1 Conducting Polymers in Biotransducer Development and Application ....................................... 12-3 Conducting Polymers as Transducer-Active Materials . Conductive Polymers as Biorecognition Layers . Conducting Polymers as Synthetic Bioreceptors
12.3
Issues and Strategies for Improved Biotransducer Analytical Performance ................................................. 12-18 Ion Transport . Stability . Adhesion . Multienzyme and Multilayer Configurations . Immobilization of Redox Mediators . Derivatization of Monomers before Conducting Polymer Electrosynthesis . Direct Electrical Modulation of Enzymes . Electroactive Polymer Hydrogels
12.4
Anthony Guiseppi-Elie Sean Brahim, and Ann M. Wilson
12.1
Biotransducer Devices and Biosensor Systems ........... 12-26 Microfabricated Array Electrodes . Electroactive Polymer Sensor Interrogation System for Conductimetric Response and Impedimetric Response . Microcantilevers . Biocompatiblity for In Vivo Sensing . Electronic Noses . Nanowire Arrays
Introduction
Biosensors are bioanalytical systems defined by their integration of appropriate levels of sample handling, biotransduction, instrumentation, and information output. Such systems have as their purpose the reporting of, through measurement or monitoring, the chemical potential energy of a targeted analyte. The biosensor measures or monitors then generates analytically useful data in an information context. Central to the biosensor system, and indeed defining over other bioanalytical systems, is the use of a biotransducer. The biotransducer is a device that converts the concentration of the target analyte into a proportionate measurable signal via an intimate combination of a biological recognition system and a physicochemical transducer. All biosensors exploit a close harmony between a selective biorecognition system and a physicochemical transducer (Figure 12.1A and Figure 12.1B). The biorecognition system is typically an enzyme, sequence of enzymes, lectin, antibody, membrane receptor protein, organelle, bacterial, plant or animal cell, or whole slice of plant or mammalian tissue. This component of the sensor is responsible for the
12-1
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Data
Sample Storage handling preparation delivery
(A)
Biotransducer
Biotransducer Interferences Foulants temperature Convection Single or multiple calibration Stds. and controls
Interrogate capture condition amplify reduce process store
Information output
Instrumentation
Biotransducer
Interferent
Signal Analyte
(B)
Biorecognition layer Physicotransducer device
FIGURE 12.1 (A) Schematic illustration of the principal components of a generic biosensor device and (B) The biotransducer component of a biosensor device is typically composed of a tandem arrangement of some biological recognition system in intimate association with a physicochemical transduction system.
selective recognition of the analyte, the generation of the biochemically derived signal monitored by the transducer, and ultimately, the specificity of the final device. In all cases, the biorecognition system is immobilized in close or intimate contact with the physicochemical transducer. The purpose of the biotransducer is to convert the biochemical signal into an electronic, optical, or magnetic signal that can be suitably processed and presented as data output. The transducer can take many forms, but the emphasis to date has been on the following electronic configurations: optoelectronic detectors used in optical biosensors, field-effect transistors (FETs), potentiometric, amperometric, voltametric, or impedimetric electrodes used in electrochemical biosensors, thermistors for thermal biosensors, and piezoelectric oscillators or surface acoustic wave devices used in mass-sensitive biosensors. Biosensors promise to provide a powerful and inexpensive alternative to conventional analytical strategies for assaying chemical species in complex matrices. Biosensors do this by their ability to discriminate the target analyte from a host of inert and potentially interfering species without the requirement for separating and subsequently identifying all the constituents of the sample. Application areas where biosensors are set to make a significant impact reach well beyond the established needs of medicine and human health, where efficient access to biochemical information has always been at a premium, to include veterinary science, environmental monitoring, food processing, bioprocess monitoring, agriculture, pharmaceuticals development, the petrochemical industry, and more recently, biodefense and other security applications [1,2].
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Polymers, both intrinsically conducting and nonconducting, have and continue to play a prominent role in biosensor device fabrication and performance. A survey of the literature on biosensors reveals that polymers have acquired a major and prominent position as key materials of construction in various biotransducer devices. An intrinsically conducting polymer may be used as a transducer-active or fieldresponsive coating or it may serve as an encapsulating material on an electrode or device surface. A nonconducting polymer may be used for the immobilization of specific bioreceptor agents on the biotransducer device and of course may be an encapsulant. This is in large part due to the versatility in the chemical and physical properties of both types of polymers, which allows them to be molecularly engineered, both in the bulk and at their surfaces, for specific biosensing applications.
12.2
Conducting Polymers in Biotransducer Development and Application
Conducting polymers have, since the very early days following their discovery, been featured in biotechnical applications [3–8]. The early work of Guiseppi-Elie and Wnek [9] sought to chemically modify the surface of polyacetylene using permanganate and peroxide solutions with the goal of introducing hydroxyl groups for enhanced wettability and for covalent immobilization of enzymes to produce hybrid structures. Considerable work has since been done involving conducting polymers that has resulted in a rapidly growing body of research, patent, and product literature on biosensors based on conducting electroactive polymers (CEPs). In general, the polymer may play an active role in transduction, in which case it is described here as transducer-active or it may play a passive role, used solely as an immobilization matrix for the biological recognition entity. Finally, unique approaches that combine biorecognition and signal transduction employ such concepts as molecular imprinting, synthetic enzymes and antibodies, covalent bioimmobilization, and synthetic receptors.
12.2.1 Conducting Polymers as Transducer-Active Materials Conducting polymers have been demonstrated to enhance desirable analytical transduction performance factors of biosensors such as the speed, sensitivity, and versatility, particularly in diagnostic applications [10–20]. The versatility of CEPs resides in the considerable flexibility in the available chemical structure. This characteristic results in the inherent ability of these materials to undergo a wide range of molecular interactions that have the potential to be ‘‘molecularly engineered’’ into the final chemical and microstructure during synthesis. Furthermore, the extent of many such interactions may be controlled in situ after synthesis, using simple electrical stimuli, thus making it possible to modulate the desired electronic and mechanical properties of the polymer. When a conducting polymer imparts an active role to a biotransducer, it may serve as a catalytic layer, a redox mediator, an ON=OFF switch, a biochemically responsive chemiresistor with a resistance value that is modulated by a targeted biochemical reaction, or may provide for molecular recognition or preconcentration of the targeted analyte. CEPs may of course be synthesized by a wide variety of methods. Discussed in detail elsewhere in this handbook, the methods of oxidative chemical synthesis and electropolymerization are most amenable to biosensor applications. The specific method of electropolymerization of a monomer from solution to a conducting polymer film directly on an electrode or device structure is amenable to direct biomolecule immobilization within the same one-step process, or correspondingly, the biomolecule may be covalently tethered to the conducting polymer backbone through complementary functional group chemistry. Because of the facile nature of electrochemically depositing conducting polymers on electrode surfaces, it is possible to control the spatial distribution of biomolecules (enzymes, antibodies, and DNA) immobilized within the polymer membrane while controlling its thickness. In the unique case of electronically wired enzymes, it is also possible to modulate the immobilized enzyme activity via the impressed current and by changing the redox state of the polymer. For the proper relay of electrons from the surface of the electrode to the enzyme’s active site, the concept of ‘‘electrical wiring’’ has been
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reported and developed by Heller and Gregg [10,21], originally as a solution to the problem of diffusion of electron-transfer mediators out of an electrode-bound membrane. Conducting polymers are likely to provide a three-dimensional electrically conducting structure for this purpose. Conducting polymers can be reversibly doped and undoped using electrochemical techniques and are accompanied by significant changes in electrical and spectroscopic properties. To the extent that these changes may be linked to bioreceptor–analyte interactions, these changes, and the extent to which they are modulated, may be used as signals for the biochemical reaction [22,23]. Several papers have documented the efficient transfer of electric charge, produced by a biochemical reaction, and conducted via a CEP layer to electronic circuits [24–27]. The electronic conductivity of conducting polymers has been demonstrated to change over several orders of magnitude in response to changes in pH and redox potential, EH, of their environment [28]. These subtle changes often give an indication of what is occurring to the biomolecules entrapped within or anchored to the polymer surface. When pH and EH are appropriately linked to an analyte-specific source of modulation, and this modulation of pH and EH are exclusively linked to changes in the electrical and optical properties of the conductive electroactive polymer membrane, then these may be a source of analytical signals that may be quantified at the electrode surface as a change in conductivity (conductimetry), a change in impedance (impedimetry), a change in potential (potentiometry), or a change in current at fixed applied potential (amperometry). Examples of such biochemical reactions include enzyme-catalyzed oxidation or reduction of the analyte that leads to the formation of reaction products (e.g., protons, hydrogen peroxide, dopant anions). Changes in the chemical potential of these reaction products are reflective of the activity of the analyte through the specific conversion rate of the immobilized enzyme, often via Michaelis–Menten kinetics. The change of chemical potential of the product then modulates the electrical and optical properties of the CEP. In all cases, these changes are the consequences of changes to the density of states; the Fermi energy of the semiconducting state that is influenced by the carrier density or mobility of polarons, and the carrier density of the metallic state (Figure 12.2). Under these circumstances, the response of the biotransducer is indirect as it arises from the influence of a chemical species that is not the analyte. Sometimes, the electronic conductivity or optical properties associated with the redox state of the CEP (doping level) may vary or be modulated by direct reaction with the analyte, resulting in direct biotransducer responses [29]. Electrochemical biotransducers based on CEPs may be characterized by three important aspects: (i) the configuration and character of the underlying substrate electrode on which the CEP membrane layer is fabricated (such as Pt, Au, glassy carbon, indium–SnO2, carbon fibers, etc.), which confers mechanical stability and governs the electrical character of the electrode–CEP interface, (ii) the chemical
Conduction band
Since the conduction band for a semiconductor is above the Fermi level, it Semiconductor at high temperature must be at an elevated temperature to have any population.
Egap At absolute zero, 0 K
Egap
Fermi level
E
2 f (E ) 1.0
Valence band
Egap
Density of energy states in conduction band of semiconductor.
3/2 r(E ) = 8 2 pm E−Egap h3 EF r(E ) The overlap of valence and conduction bands in a conductor gives population of the conduction band.
Context of Fermi level for a semiconductor
FIGURE 12.2
Egap EF
The Fermi theory of conduction in semiconducting materials.
r(E ) =
Conduction Bands
Conductor at 0 K 8
2pm3/2 E h3
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character, nano- and microstructure, and gross morphology of the material comprising the final CEP membrane itself (i.e., the conductive matrix for the immobilization of biomolecules), and (iii) the immobilized biological recognition species [30]. How these three elements come together ultimately decides the character of the CEP biotransducer. 12.2.1.1 Conductimetric Devices Conductimetric biosensors report changes in the conductivity of the CEP film while under two-point or four-point DC interrogation. The corresponding voltage drop allows calculation of the device film’s conductance, and with the device’s cell constant, the film’s conductivity. Responses are typically kinetic, although equilibrium responses (i.e., the extent of change of conductance or conductivity) may be used. Such sensors are also commonly referred to as chemiresistors. Many processes that lead to changes in carrier density or mobility, such as the interaction with electron acceptors or donors, will cause changes in conductivity of these polymers. The early work on the conductivity measurements of polyacetylene films upon doping with vapors of iodine, bromine, or AsF5, and subsequent compensation with NH3 constitutes the simplest conducting polymer gas sensors [31]. Since then, electrodeposited polypyrrole and poly-3-methylthiophene films have been doped with copper or palladium for the detection of reducing gases such as NH3, H2, and CO. MacDiarmid and Epstein [32] demonstrated that the interaction of a doped conjugated conducting polymer (e.g., D,L-camphorsulfonic acid-doped polyaniline in emeraldine base form [HCSA-doped PANI-EB]) with certain organic solvents (e.g., m-cresol) could cause a conformational transition of the polymer chain from a ‘‘compact coil’’ to an ‘‘expanded coil’’ through the so-called secondary doping process, which was found to be accompanied by a concomitant change in conductivity. Such phenomena have also been observed when CEPs were exposed to some common organic vapors such as methanol, hexane, and benzene, providing the basis for developing conjugated conducting polymer–based sensors for the detection of hydrocarbon vapors for the creation of an electronic nose platform [33]. Earlier studies of indirect conductimetric biosensors based on CEPs have been described by Wrighton and coworkers. The devices of Wrighton and coworkers [34] used an ingenious connection to a third electrode to potentiate the response of the CEP that was set between a pair of interdigitated microelectrodes (Figure 12.3). In this configuration, when the applied voltage to the microelectrodes is zero
Counterelectrode Reference electrode
Potentiostat
Conducting polymer Vgate
Vdrain
Working electrode
Idrain = s
FIGURE 12.3 Schematic representation of an electrochemical device for the determination of a polymer’s conductivity as a function of analyte and applied electrochemical potential. A cross-sectional view is shown of microelectrodes (spaced at 1–5 mm) covered with an electroactive polymer. When there is no voltage applied (or voltage offset) between the microelectrodes (Vd ¼ 0), then sweeping the applied electrochemical potential (V0) produces a cyclic voltammogram. When Vd is finite (0.01–0.05 V), sweeping the applied potential will produce a current between the microelectrodes, which is proportional to the polymer’s conductivity. (Reprinted from Swager, T.M., Acc. Chem. Res., 31, 201, 1998. With permission.)
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(Vd ¼ 0) and the electrochemical potential impressed by the potentiostat is swept over a certain range (V1>Vgate>V2), a cyclic voltammogram is obtained. When Vd is made finite (0.01–0.05 V) and Vg is again swept over a potential range, a current that is proportional to the conductivity of the CEP will be produced between the pair of microelectrodes. Guiseppi-Elie and Wilson [35] described the covalent immobilization of oxidoreductase enzymes that produced hydrogen peroxide that would oxidize the CEP and so reveal a kinetic response in its conductivity linked to substrate concentration (see Section 12.4.2 on EPSIS). Malmros and Guiseppi-Elie [36] described a redox switch, the ON=OFF of which could be linked to the extent of enzyme reaction. Swager [37] investigated several approaches to create molecular wires and stabilize receptor systems in attempts to amplify the signals from conductimetric devices employing conducting polymers. The goal was to fabricate metal ion–selective chemosensors based on low molecular weight compounds (e.g., crownether derivatives) that selectively respond to the presence of specific metal ions through an electrochemical (e.g., redox potential) or optical (e.g., absorption or fluorescence) change. Structures such as crown-ether macrocycles were incorporated into polythiophene films to introduce large resistive elements in conductive polymer chemiresistors. The calix[4]arene arrangement, also used with polythiophene films, was shown to provide a highly organized and rigid receptor structure. Furthermore, the binding of viologens to these macrocycle-containing conducting polymers lead to charge-transfer interactions between polymer and viologen, which create barriers to carrier (polaron=bipolaron) transport. Such barriers enhance the formation of resistive elements within conducting polymers. The end result of these conducting polymer hybrid structures was the fabrication of chemosensors with highly specific responses that can in principle be used to produce highly accurate CEP array devices for electronic noses or tongues. 12.2.1.2 Impedimetric Biotransducers Impedimetric biotransducers report changes in the CEP film while under a sinusoidally varying, nonperturbating, interrogating voltage (typically 50 mV p-t-p) relative to a suitable reference electrode. The ensuing current allows the transfer function to reveal the real (Z 00 ) and imaginary (Z 0 ) components (or magnitude jZj and phase u) of the impedance that change predictably as a function of the concentration of the analyte during exposure. These responses may be measured as a function of frequency over some specified frequency range, in which case it is called electrochemical impedance spectroscopy (EIS), or at a single frequency, in which case it is called electrochemical impedance (EI). The response of the biotransducer may be indirect, for example when it arises from the reaction of a product of a biorecognition reaction within the CEP membrane layer that leads to a change in carrier density (doping) or mobility (swelling). The response may be direct, for example when it arises from a change in either the real or imaginary component of the impedance that accompanies the binding of a recptand (analyte) with a bioreceptor immobilized within or at the surface of the CEP (for a recent review of EIS and EI involving CEPs applied to genosensors, see Ref. [38]). Direct impedimetric detection of biological recognition reactions leading to biotransducers was first described by Teasdale and Wallace [39]. Sargent and Sadik [40] investigated the mechanisms of antibody–antigen (Ab–Ag) interactions at conducting polypyrrole electrodes using impedance spectroscopy techniques. The theory of charge generation and transport in the heterogeneous polymeric interface was proposed to explain the current flow during Ab–Ag binding. According to this mechanism, the current obtained at the Ab-immobilized conducting polymer electrodes occurred via the following steps: (i) diffusion of ions to the electrode, (ii) charge transfer at the porous PPy–membrane interface, (iii) migration through the polymer PPy membrane, and (iv) adsorption and desorption of the Ag at the PPy–solution interface. Step (iv) was considered to be the rate-determining step and could be controlled through the appropriate choice of electrical potential. At positive potentials, Ab–Ag reactions were enhanced, and the application of negative potentials disrupted this association. These findings confirmed that the Ab–Ag interaction was largely influenced by the applied potential to the conducting polymer–modified electrode surface. Thus, the combination of Ab immobilization onto conducting polymer matrices and pulsed potential interrogation waveform was demonstrated to enable direct selective molecular recognition.
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Recently, Guiseppi-Elie et al. have described polypyrrole-based [41], and polyaniline-based biosensors [42] for the impedimetric detection of DNA hybridization. For the case of polypyrrole, single-stranded probe molecules were incorporated by electropolymerization and were found to conform to the expected ratio of one base per four pyrrole repeat units in the fully doped state. Hybridization with compliment, noncompliment, and single-base mismatched target sequences demonstrated the bioanalytical sensitivity of this device. This research investigated the use of these two conducting polymers for possible enhanced label-free electrochemical detection of DNA hybridization (Figure 12.4). DNA probes were immobilized by either physical adsorption or covalent linking to polyaniline and polypyrrolecoated electrode surfaces. When the immobilized ssDNA probe on the conducting polymer was hybridized to its redox-labeled, complementary ssDNA target, the redox moiety was then brought into close proximity to the electrode surface. Voltammetric detection (via differential pulse voltammetry) over the range 0.2–0.6 V applied to the electrode and using a modulation amplitude of 10 mV in combination with a scan rate of 100 mV=s allowed the current resulting from oxidation of the redox label to be monitored. While a full mechanistic understanding of the factors responsible for impedimetric change in conducting polymer membranes at electrode surfaces is unclear at present, there is substantial evidence that indicate that such transduction changes may be employed in truly reagentless biosensing. The use of EIS and the conducting polymer, polypyrrole, as an integrated recognition and transduction system for reagentless biosensor systems was demonstrated by Vadgama and coworkers [43]. The first system incorporated polypyrrole primarily as the immobilizing matrix for the antibody antiluteinizing hormone (anti-LH) entrapped through potentiodynamic electropolymerization from the antibody–monomer solution. The resulting antigenic biosensor showed no apparent change in Bode impedance parameters (jZj or u) upon exposure to luteinizing hormone (LH). However, following a redox cycle, a significant change emerged with regard to the peak phase angle, u. The process of redox cycling was thought to result in a realignment of the conducting polymer chains following the antibody–antigen interaction. Two charge-transfer processes were observed with these antibody-loaded polypyrrole films, which were assigned to polaronic conduction at low frequencies and electronic conduction at high frequencies. The same researchers also demonstrated a construct for DNA hybridization discrimination to differentiate single- and double-stranded DNA based on the interaction of polypyrrole with either surface-adsorbed single-stranded DNA or 10-mer oligonucleotide sequence prehybridized with its
Immobilized probe DNA
Ferrocence label
Target DNA
Fe2+
OH O= C
NH O=C
NH O=C
Electro-co polymerization N
N H
+
N H
N H
H
Electrode
N H
H N
Ferrocene
30 –100 nm
Electrode
FIGURE 12.4 Bioelectronic detection of DNA hybridization based on electrochemical detection of ferrocene label at conducting polymer–modified electrodes. (Reprinted from Lei, C., Georghe, M., and Guiseppi-Elie, A., Proc. ACS PMSE Preprints, 83, 552, 2000. With permission.)
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400
−Zim / KΩ
300
200
100
0 0
200
400
600
Zre / KΩ
FIGURE 12.5 Complex plane impedance plots for an antibody-modified electrode at a potential of 1.4 V vs. SCE and AC signal of 5 mV, in PBS (squares), without antigen injection; after antigen injection at different concentrations: circles, 10 ng=mL; positive triangles, 50 ng=mL; and inverted triangles, 100 ng=mL. (Reprinted from Ouerghi, O., Touhami, A., Jaffrezic-Renault, N., Martelet, C., Ben Ouada, H., and Cosnier, S., Bioelectrochemistry, 56, 131, 2002. With permission.)
compliment strand. Characteristically, different acquired EIS data expressed through Nyquist plots were observed following electrochemical cycling of both types of oligo-derivatized interdigitated electrode surfaces in pyrrole monomer. The significant differences between the two spectra were interpreted as due to the different reporter interactions of the conducting polymer with the double-stranded chain versus the single-stranded DNA chain. It was proposed that the polymer may be acting through a specific intercalation or groove-binding route in the case of the double-stranded helix. An electrodeposited biotinylated polypyrrole film was described as an immobilization matrix for the fabrication of impedimetric immunosensors [44]. Biotinylated antibody (antihuman IgG) was attached to the free biotin groups on the conducting polypyrrole film with avidin as a coupling reagent. Nyquist plots demonstrated an increase in the semicircle diameter (Figure 12.5) with increasing antigen concentration, especially at low frequency, which could be chosen for concentration-dependent impedance measurements. This particular immobilization method allowed a highly reproducible and stable device to be obtained. The resulting immunosensor exhibited a linear dynamic range of 10–80 ng=mL of antigen and a detection limit of 10 pg=mL. Contractor et al. [17] have fabricated impedimetric biosensors for the estimation of several analytes including glucose, urea neutral lipid=lipase, and hemoglobin=pepsin. The change in electronic conductivity of the conducting polymer component of the biosensor device was measured in response to changing the applied redox potential or microenvironmental pH of the polymer matrix. Ramanathan et al. [45] investigated the application of polyaniline-based Langmuir–Blodgett films in the fabrication of a conductimetric glucose biosensor. Conductivity biosensors based on conducting polymers were also demonstrated for penicillin and urea. 12.2.1.3 Amperometric Biotransducers Amperometric biosensor devices report the resulting current produced as a consequence of some specific electrochemical reaction, or cascade of reactions, occurring at an appropriately polarized biotransducer
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electrode surface. Amperometry is the most common approach to sensing involving electroconductive polymers. The proper functioning of amperometric biosensors is governed by the efficacy of the electron transfer between the catalytic biomolecules, usually enzymatic belonging to the oxidoreductase or dehydrogenase class, and the underlying electrode surface, most often involving redox mediation or a conducting polymer layer. To date there has been numerous reports on the use of conducting polymers in the fabrication of amperometric biotransducers. By far, the largest group of direct electron-transfer amperometric biosensors is based on coimmobilization of the enzyme in a conducting polymer matrix, most often polypyrrole and polyaniline. Table 12.1 shows cross sections of the various biosensors developed for detecting a host of analytes employing the combination of conducting polymers and amperometric transduction. Resoundingly, the majority of such biosensors make use of polypyrrole as the conducting polymer matrix in combination with the enzymatic immobilization of glucose oxidase (GOx); partly due to the medical and commercial importance of detecting glucose in biological fluids and partly because of the very robust nature of GOx that can function as a model for subsequent redox enzyme biosensor platforms. While a variety of different biological materials (whole cells, antibodies, haptens, oligonucleotides, chemical receptors) have been immobilized within or onto conducting polymers, redox enzymes are still the most widely used biomolecules considered useful from an electrochemical point of view. They are used to catalyze the conversion of the substrate between its oxidized and reduced state, following which some specie of the enzymatic reaction is detected at the electrode (Figure 12.6). As previously mentioned, oxidase enzymes (for which oxygen functions as the natural electron acceptor) and dehydrogenase enzymes (for which a solubilized physiological nicotinamide cofactor, NAD(P)þ functions as electron acceptor or as electron donor, NAD(P)H) predominate [46]. Peroxidases and pyrroloquinoline–quinone
TABLE 12.1
Biosensors Based on Conducting Polymers
Substrates or Species to Be Determined Glucose
D-Alanine
Atrazine Cholesterol Choline Glutamate Fructose Hemoglobin L-Lactate
Enzyme Glucose oxidase
Glucose dehydrogenase acid oxidase Tyrosinase Cholesterol oxidase and cholesterol esterase Choline oxidase Glutamate dehydrogenase Fructose dehydrogenase Pepsin Lactate oxidase Lactate dehydrogenase D-Amino
Polymer Polypyrrole Poly(N-methylpyrrole) Polyaniline Polyindole Polypyrrole Polypyrrole Polypyrrole Polypyrrole
Lipids Phenols Urea
Lipase Tyrosinase Urease
Substituted polypyrrole Polypyrrole Polypyrrole Polyaniline Polyphenylene diamine Polyaniline Polypyrrole–polyvinyl sulfonate Polyaniline Substituted polypyrrole Polypyrrole
Uric acid Triglycerides
Uricase Lipase
Polyaniline Polyaniline
Detection Amperometry Potentiometry Amperometry Amperometry Amperometry Amperometry Amperometry Amperometry Amperometry Amperometry Amperometry Amperometry Conductometry Amperometry Amperometry Amperometry Conductometry Amperometry Amperometry Potentiometry Conductometry Capacitance measurement Admittance measurement Amperometry Conductometry
Source: Reprinted from Gerard, M., Chaubey, A., and Malhotra, B.D., Biosens. Bioelectron., 17, 345, 2002. With permission.
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Product
Substrate
OXD (ox)
O2 + 2H+
OXD (red)
H2O2
2 e−
FIGURE 12.6 Mechanism of amperometric transduction at an electrode surface using an oxidoreductase enzyme (OXD) that catalyzes the conversion of substrate to product.
(PQQ), which are present in some dehydrogenase enzymes, have recently become of particular interest. PQQ has also been increasingly used as an artificial electron-transfer mediator with other types of redox polymers [47]. The sensing of glucose, a physiologically important analyte, has received more attention that any other of clinical significance. This is fueled by the promise of the fabrication of true closed-loop delivery systems for better and more accurate chronic diabetes management that will allow for controlled release of insulin in direct response to the output of either an implanted or noninvasive continuous glucose biosensor. Conducting polymers have been widely applied to the fabrication of glucose biosensors of varying formats. From the simplest fabrication techniques involving entrapment of an enzyme or enzymes within a conducting polymer matrix to the use of composite polymers or polymer multilayers, several amperometric glucose biosensors have been demonstrated. As earlier as the mid-1980s, the incorporation of glucose oxidase as counterion during the electrosynthesis of either polypyrrole or polyaniline was shown to be an elegant strategy to create amperometric glucose biosensors [48,49]. Sangodkar et al. [50] reported microfabricated biosensor arrays of gold for glucose, urea, and triglycerides based on incorporation into polyaniline. These conducting polymer–based biosensors were reported to accurately sense these three analytes from a sample containing a mixture of the three, thus demonstrating a test bed for an electronic tongue. The use of polyaniline as an enzyme immobilization matrix for amperometric detection of glucose was also demonstrated using a Prussian Blue– modified platinum electrode [51]. This particular biosensor construct was highly effective in countering the influence of common electro-oxidizable interferents. The entrapment of glucose oxidase into polypyrrole films followed by extensive polarization to produce nonconducting versions of the films has been a successful procedure for imparting enhanced selectivity to amperometric glucose biosensor systems [52]. It is thought that this process results in disruption of the conjugation at several points along the polymer backbone, which results in the loss of electronic conductivity, with the successive generation of electronegative functionalities. Such groups on the polymer backbone repel the negatively charged electrooxidizable interferents at physiological pH (Figure 12.7). Similar anti-interferent properties were endowed to glucose-biosensing layers created by first overoxidation of PPy films followed by gel entrapment of glucose oxidase in bovine serum albumin (BSA)=glutaraldehyde deposited over this oxidized PPy inner layer [53]. The combination of glucose oxidase immobilized in overoxidized polypyrrole and microdialysis was used for continuous subcutaneous detection of glucose in a rabbit [54]. The use of two-enzyme systems in conducting polymers to enhance selectivity of amperometric biosensors has also been demonstrated. Shin et al. [55] fabricated glucose biosensors by entrapment of glucose oxidase in polypyrrole followed by a top layer containing only horseradish peroxidase (HRP). HRP is well known to catalyze the oxidation of several interferents in the presence of its natural substrate, H2O2. The resulting preoxidized interferents will then contribute no ‘‘false’’ amperometric signals when impinged on the underlying polarized metal electrode (Figure 12.8). This strategy proved to be somewhat effective only at a high enough concentration of diffusing-away H2O2, generated at relatively high glucose concentrations. Tian and Zhu [56] entrapped glucose oxidase in polypyrrole that was now the outer layer covering an HRP-modified sol–gel-derived mediated ceramic carbon electrode.
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A
Platinum electrode
Polypyrrole Anionic interferent
−
Ox
−
Ox
Ox
−
−
I−
I−
Ox
I−
Bulk medium Substrate
FIGURE 12.7 Schematic showing the mechanism of anion interferent shielding using extensively oxidized polypyrrole. The prolonged oxidation of the conducting polymer is thought to result in loss of inherent conductivity concomitant with the generation of electronegative functionalities on the polymer backbone.
The electron transfer mediator, ferrocene carboxylic acid, was incorporated into the system, which allowed the electrode to detect glucose at the low potential of þ0.16 V. Both enzymes showed favorable retention of activities in this polymer configuration. Redox polymers have been combined with conducting polymers to fabricate amperometric biosensors for glucose. Rohde et al. [57] described a reagentless, flow-through glucose biosensor by codeposition of glucose oxidase with glutaraldehyde and the osmium-based redox polymer synthesized by Heller
GOD HRP Glucose GOD H2O2
H2O2
Interferent HRP
GOD
HRP
Oxidized interferent
GOD
Pt electrode
Underlying layer
Interferent + Glucose + H2O2
HRP
Overlying layer
Bulk solution
Oxidized interferent + Glucose + 2H2O
FIGURE 12.8 Mechanism of preoxidation of potential interfering species using horseradish peroxidase in the presence of its substrate hydrogen peroxide. (Reprinted from Shin, M.-C., Yoon, H.C., and Kim, H.-S., Anal. Chim. Acta, 329, 223, 1996. With permission.)
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[Os (bipyridyl)2(poly-4-vinylpyridine)10Cl]Cl, on an underlying platinum electrode, followed by an outer layer of electropolymerized polypyrrole containing glucose oxidase. This biosensor was capable of detecting glucose linearly up to 50 mM, with a reported sensitivity of 12.2 nA=mM. Glucose oxidase is an enzyme well known for its robustness, and this fact has been taken advantage of in fabricating amperometric glucose biosensors containing the enzyme covalently linked to conducting polymers. Schuhmann [58] described the preparation of such biosensors whereby glucose oxidase was covalently bound to the outer surface of functionalized polypyrroles for the detection of glucose in flowinjection manifolds [58]. These amperometric sensors showed high sensitivity and were effective in screening out electro-oxidizable interferents from solutions containing the analyte along with interfering compounds. The use of a viologen mediator for covalent tethering of glucose oxidase onto polypyrrole was demonstrated by Liu et al. [59]. The viologen (N-(2-carboxyl-ethyl)-N 0 -(4-vinylbenzyl)-4,40 -bipyridinium dichloride or CVV) was the first graft polymerized onto polypyrrole, which then served as covalent anchor to the enzyme through the available carboxylic acid functionalities. The resulting glucose biosensor was capable of linearly detecting glucose up to 20 mM, with 40% loss in enzyme activity after storage in buffer for 10 days. In another study, glucose oxidase was coupled to a strong polyanion, poly(2-acrylamido-2-methylpropane sulfonic acid) (AMPS), via polyethylene glycol (PEG) spacers to effectively and reproducibly immobilize the enzyme within a polypyrrole matrix onto a Pt electrode surface [60]. PEGs with four different chain lengths (1000, 2000, 3000, and 4000) were used as spacers to study the spacer length effect on enzyme immobilization and electrode function. After conjugation, more than 90% of the immobilized enzyme bioactivity was preserved and the bioactivity of the conjugated enzyme increased with longer PEG spacers. The amperometric response of the conducting polymer–based biosensor was demonstrated to linearly detect glucose up to 20 mM with a reported sensitivity ranging between 180 and 270 nA=mM=cm. Wang and Musameh [61] demonstrated the one-step incorporation of both carbon nanotubes (CNTs) and glucose oxidase into electropolymerized polypyrrole for the fabrication of amperometric biosensors with enhanced sensitivity and selectivity. The nanotube dopant retains its electrocatalytic activity towards enzymatically generated hydrogen peroxide, with a linear dynamic range up to 50 mM glucose and sensitivity of 2.33 nA=mM. Such simultaneous immobilization of both nanotubes and glucose oxidase imparts electrocatalytic and biocatalytic properties onto amperometric transducers and offers the promise of fabricating glucose nanosensors based on recently developed polypyrrole=CNT nanowires [62]. Increasing the effective surface area or roughness of the underlying platinum transducer surface of amperometric biosensors has also been attempted to promote enhanced biosensor sensitivity [63]. Electrodeposited Pt black and templated Pt nanowire brushes were explored as potential Pt deposits. These deposits resulted in orders of magnitude increase in sensitivity, attributed to the increased loading of glucose oxidase onto the larger electrode surface areas and to the higher hydrogen peroxide electrooxidation efficiency. Similar improved sensitivity results were observed by the use of boron-doped diamond microfiber (BDDMF) electrodes modified with platinum nanoparticles. Polypyrrole films were electrochemically grown over the Pt=BDDMF electrodes and subsequently overoxidized by the application of þ0.7 V (versus Ag=AgCl) for 1 h. These polymer films were then made glucose-responsive by immobilization of glucose oxidase through BSA–glutaraldehyde cross-linking over the oxidized PPy. The biosensor exhibited linearity to glucose up to 70 mM and was effective in screening common endogenous interferents such as ascorbate. In vivo biosensors for glucose based on derivatized conducting polymers were also pursued. Yasuzawa et al. [64] synthesized pyrrole monomers that were derivatized with the phosphotidylcholine group to function as hemocompatible biosensors. This moiety is the principal component of the outer leaflet of red blood cell membranes. Subsequent electropolymerization of the pyrrole derivatives in the presence of glucose oxidase formed the outer bioactive coating of a glucose biosensor that also contained an inner membrane of Nafion. The combination of these membranes resulted in enhanced selectivity and retention of enzyme activity. Another clinically important analyte that has received much attention for continuous monitoring via amperometric transduction is lactate. Interstitial levels of the L-isomeric form have been demonstrated
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to be the reflective of systemic levels that correlate with the severity of traumatic injury, including hemorrhage [65,66]. Ru¨del et al. [67] designed novel enzyme reactors with immobilized lactate oxidase in polypyrrole by electrochemically depositing enzyme–polymer layers on pieces of graphite felt that were subsequently mounted into the column of a flow system. With this system, a linear range up to 1.0 mM lactate and detection limit lower than 50 mM was achieved. Chaubey et al. [68] demonstrated the coimmobilization of lactate oxidase and lactate dehydrogenase by physical adsorption onto electrochemically prepared polyaniline films. The resulting bienzyme electrodes were shown to provide signal amplification through substrate recycling with a linear dynamic response range from 0.1 to 1.0 mM lactate and with a detection limit of 1.0 105 M. The same group previously fabricated amperometric lactate biosensors that employed lactate dehydrogenase immobilized via glutaraldehyde cross-linking onto electrochemically grown polypyrrole polyvinylsulfonate composite films [69]. Analytical performance characteristics of linear range 0.5 to 6.0 mM, detection limit of 1.0 104 M, and response time of 40 s were achieved. The research group of Guiseppi-Elie et al. [70] designed an amperometric biochip for monitoring both lactate and glucose in an effort, among others, to investigate the intimate physiologic relationship between these two paramount metabolites within a muscle bed during hemorrhage. The biorecognition layer of the biochip consisted of a composite hydrogel membrane containing a unique bifunctional polypyrrole component immobilized over a novel platinum microdiscarray sensor design. The pyrrole monomer was synthesized to bear an UV-polymerizable methacryloyl terminus for covalent cross-linking into the hydrogel, and an oxidatively polymerizable pyrrolyl terminus, for propagating the conducting polymer component throughout the network. The amperometric biochip was capable of detecting glucose linearly over the range 0.1 to 13.0 mM and up to 90.0 mM for lactate. 12.2.1.4 Potentiometric Biotransducers In comparison to the overwhelming quantity of literature on the use of conducting polymers combined with amperometric detection, the coupling of conducting polymers with potentiometric biosensing has been considerably less studied. While there have been several reports on the use of conducting polymers in the fabrication of potentiometric ion chemical sensors [71], there are only a handful of reports that describe the fabrication of potentiometric biosensors using either polypyrrole or polyaniline as the immobilization matrix. Adeloju and Moline [72] demonstrated the fabrication of ultrathin films of polypyrrole containing the enzyme glucose oxidase on platinum disc electrodes. These 55 nm thick membranes were electrosynthesized from electrolyte-free pyrrole solution containing the enzyme via galvanostatic means. The optimized conditions for growing the ultrathin films were 0.1 M pyrrole monomer (deaerated), 55–110 U=mL of enzyme, an applied current density of 0.05 mA=cm2, and an electrical charge of 25 mC=cm2. The addition of an outer thin membrane of PPy–Cl to create a bilayer biosensor configuration effectively suppressed the potentiometric interference caused by ascorbate. The potentiometric glucose biosensor was demonstrated to detect glucose linearly over the range 0.06–10 mM. Several other glucose biosensors employing conducting polymers of polypyrrole and polyaniline and potentiometric transduction have been reported. Glucose oxidase was entrapped within a conducting thin film of polypyrrole that was doped with chloride [73]. It was found that the thinner the growth of the bioactive film, the greater the resulting sensitivity of the potentiometric biosensor. The main drawback with this biosensor, however, was the small linear dynamic response range observed. In another study, the researchers entrapped glucose oxidase within a thin film of polyaniline that was coated with an external layer of Nafion [74]. Like the previously described potentiometric PPy–Cl glucose biosensor, the NafionjPAn-GOx biosensor also displayed a very limited linear dynamic response, 1–3 mM glucose. The potentiometric sensing of urea was also demonstrated by entrapping the enzyme urease within a polypyrrole membrane electrosynthesized on a platinum electrode [75]. The bioactive sensing membrane was synthesized by potentiostatic electropolymerization of pyrrole monomer in a solution containing the urease enzyme and a nucleophilic electrolyte such as NaOH, NaCO3, or NaHCO3. This biosensor exhibited a good Nernstian response to urea, with a slope of 31.8 mV=decade over an urea concentration range of 1.0104 mol=dm3–0.3 mol=dm3.
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12.2.1.5 Field-Effect Transistors The field-effect transistor is a potentiometric device that relies on the application and modulation of an electric field to regulate the shape and hence the conductivity of a ‘‘channel’’ in a semiconducting material. The field effect controls the current through a gate electrode thereby opening the possibility of a transistor action without requiring the existence of p–n junctions. This phenomenon not only improves the device characteristics but also provides a useful tool to study semiconductor and surface states. Koezuka et al. [76] first demonstrated the use of this principle to fabricate conducting polymer– based transistors. They demonstrated that it was possible to control the current flowing between source and drain through the gate (Figure 12.9). Since the pioneering work of Koezuka, many researchers have applied this principle to fabricate FET devices incorporating different conducting polymers [77–79]. Electrically conducting polymers such as highly doped polyaniline [80], polythiophene [76], and poly(3,4-ethylenedioxythiophene) (PEDOT) [81,82] have been used as not only the active channel but also as electrodes of the all-polymer-based FETs whereas various polymers such as poly(vinyl phenol) [83], polyimide [84], poly(vinyl alcohol) (PVA) [85], and optical adhesive [86] were used as the dielectric layer replacing conventional silicon dioxide or silicon nitride. One of the important parameters deciding the performance of FET devices is the mobility, and this has probably been an area of concern regarding the use of conducting polymers in these devices. The charge carrier mobilities in these polymer-based FETs are typically around 105 cm2=V s depending on the applied voltage and the nature of the gate insulator. These values are about 4 to 5 orders of magnitude lower than corresponding mobilities registered with the use of inorganic semiconductor devices. However, the use of highly stacked p-conjugated oligomers and polymers such as sexithiophene, antradithiophene, and regioregular poly(3-methylthiophene) results in high field-effect mobilities of 102 to 1 cm2=V s in their neutral state [87–89]. These values are much greater than the mobilities for amorphous p-conjugated polymers such as polyacetylene, polythiophene, and poly(3-methylthiophene) in their neutral or lightly doped states [88,90]. Apparent mobilities of positive charge carriers were found to increase by electrochemical or chemical doping [91]. While the performances of conducting polymer–based FETs are quite encouraging and have spawned abundant applications of these devices in logic circuits or active matrix emissive displays [92,93], the slower response and limited lifetime of conducting polymer–based FET devices compared to conventional Si technology at present limit the possibility of a full transition to all polymer FETs.
VG
VG Counterelectrode Reference electrode
BPS
BPS
VG > VT ≡ ON VG > VT ≡ OFF A VD ID < IT ≡ OFF
A VD ID < IT ≡ OFF
FIGURE 12.9 Schematic illustration of chemical and biological sensor configurations involving the use of conducting electroactive polymers in redox switches. (Reprinted from Guiseppi-Elie, A., Wallace, G.G., and Matsue, T., Handbook of Conducting Polymers, Marcel Dekker, New York, 1998. With permission.)
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12.2.2 Conductive Polymers as Biorecognition Layers The popularity of using CEPs in biotransducer motifs is partly due to the availability of a spectrum of immobilization techniques for bioreceptor (molecule through cell) localization, each of which is compatible with maintaining bioactivity, faithful presentation of the bioreceptor to the analyte, and controlled thickness of the biorecognition layer. Bioreceptor immobilization may be performed via any one of the following techniques: (i) adsorption (physical, chemical, and electrostatic), (ii) adsorption followed by cross-linking, (iii) ion-exchange doping, (iv) covalent linking, and (v) entrapment. 12.2.2.1 Adsorption This process involves the electrophoretic migration and incorporation of charged (negative) biomolecules, which accumulate in the vicinity of the working electrode and effect a precipitation onto the conducting polymer surface. The adsorptive process occurs in the polymer–solution interface due to resulting electrostatic interactions between the polycationic matrix of the oxidized polymer and the net negative enzyme charge. To ensure the biomolecule remains with a net negative charge, electropolymerization is performed in solution with pH greater than the isoelectric point of the enzyme. The loading capacity with this immobilization technique is restricted to very small quantities (one monolayer on the polymer surface). In spite of this, the technique has been successfully used in the fabrication of biosensors for glucose [94], cholesterol [95], and pyruvate [96]. 12.2.2.2 Adsorption Followed by Cross-Linking This method presents a simple yet compromised immobilization technique for localizing biomolecules on conducting polymers using BSA followed by glutaraldehyde in prescribed amounts. Generally, the technique results in biosensors with relatively rapid response times but decaying bioactivity (stability) due to the fact that the biomolecule is not shielded from the bulk solution and also may be partly denatured by the cross-linking process. Several short-term biosensors have been fabricated by this convenient technique, including amperometric biosensors for salicylate [97], lactate [98], and tyrosine [99]. 12.2.2.3 Ion-Exchange Doping This immobilization technique makes use of the reversible redox states of conducting polymers for incorporating, or doping, the biomolecule into the growing polymer film during potentiodynamic electropolymerization. In addition, the mobility of accompanying counterions from solution as well as the ion-exchange properties of the conducting film governs the immobilization characteristics by this method. The principle is that the growing polymer will have a net negative charge in the medium at a pH value higher than the isoelectric point of the enzyme, which facilitates doping of the biomolecule into the film during the oxidative cycles. This technique has been particularly successful with polyaniline because its electrosynthesis involves working with strongly acidic media, which promotes enzyme denaturation. Several enzymes have been reported to be immobilized by this process, including glucose oxidase [100], peroxidase [101], cholesterol oxidase [102], galactose oxidase [103], xanthine oxidase [104], and uricase [105]. 12.2.2.4 Covalent Linking This process can be subdivided into two categories depending on the moiety that is functionalized: (a) covalent linking of biomolecules to previously functionalized conducting polymers (two-step process) or (b) electrosynthesis of polymers from monomers that have been previously derivatized with the biomolecule of interest (one-step process). The former permits selection of the optimum reaction conditions for each of the two steps involved, with immobilization of the biomolecule only occurring on the outer surface of the conducting polymer. Covalent integration of enzymes in monomers before electropolymerization allows larger loading capacities. Both procedures have been applied to the fabrication of glucose biosensors [106–108].
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12.2.2.5 Entrapment This one-step technique is the simplest method of biomolecule immobilization to conducting polymers. The enzyme or biomolecule is dissolved in a solution containing the monomer; care should be taken to ensure that the pH value is only slightly acidic and will not result in denaturation. The conducting polymer film is then grown electrochemically and incorporates the homogenously distributed enzyme molecules. The obvious advantages of this process are its simplicity and reproducibility. In addition, the technique lends itself to the fabrication of multianalyte biosensors through the incorporation of coenzymes or other enzymes by adding them to the monomer solution. Not surprisingly, the majority of research using this technique has been demonstrated for immobilizing glucose oxidase in conducting films of polypyrrole or polyaniline [109–112]. A host of other enzymes have also been entrapped in various conducting polymers such as cholesterol oxidase [113], peroxidase [114], lactate oxidase [100], ascorbate oxidase [115], sulfite oxidase [116], polyphenol oxidase [117], invertase [118], uricase [119], tyrosinase [120], NADH dehydrogenase [121], as well as other enzymes that require coentrapment of their respective coenzyme, such as pyruvate oxidase þ FAD [122], and glutamate dehydrogenase þ NAPD [123].
12.2.3 Conducting Polymers as Synthetic Bioreceptors 12.2.3.1 Biospecificity through Molecular Biorecognition The most convenient, and hence most popular, approach to confer biospecificity to CEP layers is via incorporation of naturally derived bioactive molecules. For this purpose, the methods outlined in the previous section have been widely utilized. However, the use of recombinants specifically engineered for advantage in immobilization within CEPs is currently developed. 12.2.3.2 Biospecificity through Biomolecular Imprinting The technique of molecular imprinting (to create molecularly imprinted polymers, MIPs) has been successfully applied to conducting polymers. The principle involves the copolymerization of two functional monomers with cross-linking around an analyte molecule, followed by the subsequent removal of the analyte template. The resulting MIP has analyte specificity conferred to it through the cavity created at the binding sites. This specificity toward the analyte may be achieved through one or a combination of complementary polymer characteristics including molecular size, shape, and chemical functionality. The combination of conducting polymer with nonconducting polymer segments to form MIPs used for the detection of doping anions and their analogs has been a popular strategy [124]. By removal of the doping molecule, the cavity left behind in the MIP acts as the site of recognition. Polypyrrole-based MIPs have been generated in this fashion through the overoxidation of polypyrrole, with the resulting MIP specific to various biomolecules that were incorporated as the doping molecules, such as Lglutamate [125], amino acids [126], and glucose [124]. The latter system involved the incorporation of electrosynthesized poly(o-phenylenediamine) that was imprinted by glucose, and is the first reported case of an electropolymerized MIP system imprinted with a neutral template. Conducting polymers by themselves, in particular polypyrrole, have also been widely used as the functional units for molecular imprinting. The polypyrrole film possesses a positive charge allowing the formation of imprints with anionic template molecules that subsequently exhibit good selectivities. The resulting MIPs have found applications in enantiomeric separations [127], in assays, and in the construction of sensors [128,129]. In the latter case, sensors for a host of enantiomeric amino acid species have been developed, including electrochemical PPy tyrosine sensors and surface plasmon PPy sensors for the detection of ochratoxin A [130]. Potentiometric sensors for the detection of nitrate were fabricated by electropolymerizing pyrrole onto glassy carbon electrodes in the presence of NaNO3 [131]. Amperometric sensors for morphine were demonstrated by molecularly imprinting the molecule into electrosynthesized PEDOT [132]. The sensor was capable of discriminating between morphine and its interfering analog codeine.
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12.2.3.3 Synthetic Enzymes and Antibodies as Biorecognition Entities Antibodies to rhesus (Rh) antigens are important indicators in screening hemolytic disease of the newborn (HDN) and autoimmune hemolytic anemia (AIHA). Identification of the Rh antibodies formed by immune stimulation is also essential in order to maximize the in vivo survival time of transfused erythrocytes. The present technique that is widely used for the detection of these antibodies is via agglutination-based assays that are time-consuming. A prototype of an immunobiosensor for detecting antibodies recognizing the rhesus blood group antigen, Rh (D), was developed by Wallace and coworkers [133]. Human erythrocytes were incorporated into a conducting polypyrrole–polyelectrolyte matrix. The process was followed by using oximetry and light microscopy to demonstrate the integrity of the erythrocytes in the polymerization solution and in the polymer matrix; cyclic voltammetry and resistometry for electrochemical characterization of the polymer and then agglutination, ELISA techniques, and cyclic resistometry for analysis of the immunoresponse from antigen–antibody binding. Antigen–antibody binding could be detected qualitatively by using resistometry while cycling the polymer between þ0.35 and 0.7 V (versus Ag=AgCl). A characteristic cyclic change in resistance (a resistogram) was recorded. After the addition of anti-Rh (D) antibody (250 mg=mL), the change in resistance during the resistogram decreased by 1.1 V (p < 0.0008) in polymers containing Rh (D)-positive erythrocytes, whereas polymers without erythrocytes showed no significant change. In another study, Cosnier and coworkers [134] demonstrated the synthesis and use of a novel biotinlabeled ruthenium (II) tris(bipyridyl) complex functionalized by four pyrrole groups (Figure 12.10). The complex was electropolymerized to yield a polypyrrolic film that exhibited complementary binding to biotinylated enzymes and antibodies through the natural biotin–avidin complementary-binding route. The researchers showed the ability of the biotinylated photosensitive polymer to immobilize biotinylated glucose oxidase and hence function as an amperometric biosensor for glucose. The possibility to exploit the photoelectrochemical properties of this biotinylated polymer film for the transduction of surface molecular recognition without labeling of the target was examined with the detection of anti-cholera toxin antibody as model system. For this purpose, the biotinylated film was applied to the conjugation of avidin and subsequent binding of cholera toxin B subunit biotin-labeled via avidin–biotin bridges. The analyte, anticholera toxin antibody, thereafter bound, by immunoreaction, the corresponding immobilized cholera toxin B subunit epitopes, the modified electrode thus constituting a potential immunosensor. Resulting significant changes in photocurrent magnitude upon binding of the anti-cholera toxin confirmed the generation of signal in the immnunosensor was due to specific antigen–antibody interactions.
O N
HN
H N S
O
N
O N I +N N R u2 N N N
O O
S
N HN
NH O
N
FIGURE 12.10 Structure of the tris(bipyridyl)ruthenium(II) complex. (Reprinted from Blanco-Lopez, M.C., LoboCastanon, M.J., Miranda-Ordieres, A.J., and Tunon-Blanco, P., Trends Anal. Chem., 23, 36, 2004. With permission.)
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The group of Guiseppi-Elie [135] demonstrated the design, fabrication, and operation of a polypyrrole-based conductimetric biosensor for glucose [135]. The transducer consisted of an electropolymerized and covalently adhered thin electroconductive multilayered polypyrrole derivative membrane on an interdigitated microsensor electrode array followed by the addition of a third layer that conferred biospecificity via covalently immobilized biotin that imparted the capability to perform biotin–avidin= streptavidin-binding bioassays. Using the Electroconductive Polymer Sensor Interrogation System (EPSIS) method, these transducers were studied for their response to glucose. The principle of operation of the biosensor was based on two phenomena: (i) the very strong binding (affinity) between biotin and streptavidin, which immobilizes the GOx enzyme to the PPy–IME substrate and confers glucose sensitivity and (ii) the oxidation of polypyrrole by enzymatically generated H2O2, shown coupled in the following equations: GOx
b-D-Glucose þ O2 ! Gluconic acid þ H2 O2 H2 O2 þ 2Hþ Cl þ PPy ! 2H2 O þ (PPy þþ 2Cl ) The H2O2 causes oxidation of PPy resulting in a concomitant and stoichiometric change in inherent conductivity, which was recorded by the instrument. Further details can be found in Section 12.4.
12.3
Issues and Strategies for Improved Biotransducer Analytical Performance
12.3.1 Ion Transport With conducting polymers acting as three-dimensional matrices for the immobilization of enzymes where reactants are converted to the products, several biosensors have been developed over the past two decades. One of the salient characteristics of conducting polymers that has been explored particularly for separation processes is the somewhat adjustable ion-transport feature of these materials. It is known that application of an appropriate electrochemical stimulus (applied potential) results in an inherent change to the redox state of the polymer, making it either more ionically conducting (oxidation state) or less ionically conducting (reduction state). Freshly prepared conducting polymers are formed in a conductive, oxidized state with the overall charge of the polymer remaining electrically neutral (positive charge on polymer backbone balanced by negative counteranion). Upon reduction, the conducting polymer, now electrically insulating, acquires a net negative charge (neutral polymer backbone with negative counteranion). In order to maintain electrical neutrality, the dopant counteranion must either be removed by egression or alternatively the negative charge can be countered by ingression of appropriate cations. Subsequent reoxidation of the conducting polymer either reincorporates the previously expelled anion or expels the incorporated cations to achieve electroneutrality. The ability to selectively and controllably gate the voltages that cause a particular redox state to predominate in the polymer and hence facilitate or retard ion transport has been exploited by several research groups. Guiseppi-Elie et al. [136] demonstrated the controlled electrorelease of Ca2þ ions from interpenetrating networks of inherently conductive polymers, polypyrrole and polyaniline, formed within water-swellable, electrode-supported, or free-standing poly(2-hydroxyethyl methacrylate) (p(HEMA))-based hydrogels. Holding the polymer composite at initial reducing potential, there was an observable ingress and exchange of Ca2þ ions with the native Kþ ions within the gels. Sweeping anodically to oxidize the conducting polymer component resulted in exclusion of further Ca2þ ions from entering the swollen gel network. Resweeping cathodically again resulted in an influx of Ca2þ ions to maintain charge neutrality. Studies conducted by Ralph and coworkers [137] demonstrated the ability to electrochemically induce movement of both electroinactive as well as electroactive metal ions from one solution to another across free standing and composite polypyrrole and polyaniline membranes. These studies revealed that varying the size of the dopant resulted in significant variations in the flux of
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metal ions across the conducting polymers, with generally larger membrane selectivities towards metal ions observed with polymers incorporating smaller dopant molecules. The same group investigated the metal-ion transport properties of PPy membranes incorporating chelating agents 8-hydroxyquinoline-5sulfonic acid and 2,9-dimethyl-4,7-diphenyl-1,10-phenanthrolinedisufonic acid. They reported that PPy membranes containing the former dopant were very permeable to several alkali metal, alkaline earth metal, and transition metal ions under the influence of an applied electric potential. In contrast, PPy membranes doped with the latter specie was far more selective, allowing only a small range of metal ions to diffuse through the polymer. These observations strongly suggest that metal–dopant interactions within these conducting polymers are also important factors for determining metal-ion permeability. From such results therefore, it is not difficult to envisage the incorporation of these compositeconducting polymer films or membranes into active sensing regions of biosensor devices whereby through subtle modification of applied electric potential, the sensing layer can be made selective to facilitate entry of desired analytes or ions while simultaneously screen out undesirable interferences.
12.3.2 Stability Quite often there has been concern regarding the retained stability and activity of bioactive moieties immobilized within conducting polymers. However, the issue of stability as it relates to the inherent integrity of the polymer itself is also of equal importance. In this regard, two types of stability are of importance for conductive polymers attached to surfaces: (i) chemical stability during long-term usage and (ii) electrochemical stability during the use of electrochemical properties of the material. Two features contribute to the overall stability of electrodeposited polymer films: (i) the inherent stability of the conductive material and (ii) the stable deposition of the electroactive film on the substrate surface. One approach that has been reported to improve the stability of conductive polymer films is the technique of electrosynthesis under an applied centrifugal field. Atobe et al. [138] and Eftekhari [139] have separately reported on the use of centrifugal force for the preparation of highly stable films of polypyrrole and polyaniline. A modified centrifuge system engineered to contain a three-electrode electrochemical cell setup was assembled and conducting polymer grown on various metal substrates subjected to different centrifugal forces. With this setup, Atobe et al. [138] showed that the rate of polymerization, chemical and physical properties, and morphological structures of polyaniline films were significantly and aniostropically affected by the impact of centrifugal acceleration forces over the range 1–300 g. Generally, the rate of polymerization was observed to initially decrease at 290 g and then increase at >300 g compared to the polymerization rate at 1 g. The surface structure of polyaniline film grown under 300 g force appeared more uniformly deposited with a finer network compared to polyaniline grown at 1 g whereas the morphology of conducting polymer grown under 290 g force showed more granular deposits that were not networked. Similar studies conducted on polypyrrole by Eftekhari [139] produced films that exhibited higher electronic conductivities when grown under higher centrifugal forces and confirmed that conducting polymers grown under the application of strong centrifugal fields possess enhanced stability. Biosensors fabricated with such films will undoubtedly contribute to the overall stability of these devices.
12.3.3 Adhesion Conductive polymers have widespread use in the fabrication of electronic devices such as sensors and actuators, and have recently been employed in the construction of artificial muscles. All these applications typically involve electrodepositing the polymer onto an underlying, electrically conducting substrate to form a bilayer device. The subsequent redox functioning of the conductive polymer is associated with, in aqueous or organic media, the migration of ions into and out of the film, with concomitant dynamic changes to the dimensions of the polymer. This repeated redox cycling generates interfacial shear stress, which typically culminates in delamination of the polymer film from the conducting substrate, since the bonding between the polymer film and the underlying conducting
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substrate is primarily due to hydrophobic interactions. The combined problems of poor adhesion and subsequent delamination of conductive polymer films have been the major technological limitation to the development of viable chemical and biological sensor devices. Several efforts have been investigated in attempts to improve the adhesion of conductive polymers onto conducting surfaces for long-term applications. A series of thiol-modified pyrroles were synthesized and evaluated as a monolayer to bridge the electrode surface and the top layer of polypyrrole [140–143]. However, even though the thiol portion was covalently bonded to the gold surface, there was not much convincing evidence for improved adhesion between the distal pyrrolyl terminus and the polypyrrole layer. In fact, it was suggested that slight improvements to adhesion was instead due to the electrode surface heterogeneity (roughness) or coupling of the conductive polymer with decomposition products of the monolayer. The former possibility was verified in another study [144] using chemically oxidized titanium as the substrate for growth of polypyrrole. The conducting polymer survived as much as 6000 redox cycles without any signs of delamination. Atomic Force Microscopy (AFM) analyses confirmed that the chemically oxidized Ti surface was three times rougher than untreated Ti surfaces and this feature was presumed to account for the marked improvement in polymer adhesion to electrode surface. In the early work of Guiseppi-Elie et al. [145], the use of terminally active silanes and heterobifunctional cross-linkers to promote specific adhesion of electropolymerized polymers to surfaces was investigated. Polypyrrole thin films were electrodeposited onto interdigitated microsensor electrode arrays composed of silica-based substrates containing magnetron-sputtered gold (IME-1550-M-P). To improve the adhesion between the conductive polymer and the borosilicate glass, derivatization was achieved by chemical modification with 3-aminopropyltrimethoxysilane followed by direct linking of the primary amine to the carboxylic acid functionality of 3-(1-pyrrolyl)propionic acid using the heterobifunctional linker N-hydroxysulfosuccinimide enhanced 1,3-diisopropylcarbodiimide in aqueous solution. Films grown in this manner were characterized by the time to adhesive failure using the adhesive tape test following immersion in PBKCl 7.2 buffer or in an inert dry environment. The polypyrrole films grown over the chemically treated surfaces displayed longest time to adhesion failure (>235 d) under both wet and dry conditions, superior to polymer films that were electropolymerized on untreated electrode surfaces (3 d), or organosilane only derivatized surfaces (5–27 d). Another interesting approach, albeit an indirect technique, by the group of Guiseppi-Elie et al. [70] involves the use of a bifunctional pyrrole monomer electropolymerized within a p(HEMA)-based hydrogel that was covalently anchored to the underlying borosilicate glass substrate portions of an electrode. The bifunctional monomer, 2-methacryloyloxyethyl pyrrolyl butyrate, was synthesized in our laboratories by the esterification of 1H-pyrrole-3-butyric acid (also synthesized in our laboratories) with 2-hydroxyethyl methacrylate (commercially available) as shown in Scheme 12.1. This bifunctional monomer has a UV-polymerizable methacroyl terminus and an oxidatively polymerizable pyrrolyl terminus and may serve as a cross-linker between p(HEMA) chains and polypyrrole chains. Polymerization of this derivative covalently secures the electroactive polypyrrole component to the hydrogel network, which in turn is covalently attached to the previously functionalized and derivatized substrate. Such polymer conetworks were demonstrated to swell but not delaminate from the underlying surface when immersed in aqueous solutions for extensive periods [146].
12.3.4 Multienzyme and Multilayer Configurations There are several advantages conferred to biosensor systems by incorporating a multitude of enzymes into the sensing layers. For example, the coimmobilization of peroxidase with different oxidases in conducting polymers has been demonstrated to reduce the applied potential for electrochemical oxidation of enzymatically generated hydrogen peroxide [147]. A bienzyme electrode for the detection of total cholesterol was fabricated by incorporating in situ cholesterol esterase and cholesterol oxidase in polypyrrole films during electropolymerization [148]. For the detection of creatinine, Yamato et al. [149] coimmobilized three enzymes, creatininase, creatinase, and sarcosine oxidase, into a polypyrrole matrix. The product of each enzymatic reaction, starting with creatininase, functioned as the substrate for the subsequent catalysis reactions.
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CH3 CH2
C C O
H O C CH2 CH2 CH2
+
O
O CH2
N H
1H-Pyrrole-3-butyric acid
CH2 OH 2-Hydroxyethyl methacrylate DCC
DMAP RT, 24 h
CH3 H2C C C O O CH2 CH2 O C CH2 CH2 CH2 O
N H
2-methacryloyloxyethyl 3-pyrrolylbutyrate
SCHEME 12.1 Synthesis of the bifunctional monomer, 2-methacryloyloxyethyl-3-pyrrolyl butyrate, via esterification reaction between 2-hydroxyethyl methacrylate and 1H-pyrrole-3-butyric acid. (Brahim, S., A.M. Wilson, D. Narinesingh, E. Iwuoha, and A. Guiseppi-Elie. 2003. Chemical and biological sensors based on electrochemical election using conducting electroactive polymers. Microchim Acta 143 (2–3):123.
Polymer multilayers offer an increased dimensionality to single-layer biosensors and provide a number of alternatives for improving sensitivity and selectivity. One such alternative draws on the advantage of increased perm selective properties of several polymers [150–152]. Another possibility affords a reduction in the working overpotential by incorporating bilayer systems of a conducting polymer and a redox polymer layer [153]. Enhanced selectivity towards the desired analyte can be accomplished by combining enzymatic systems localized in different layers in the polymeric matrix. One common approach is to use an outer layer containing immobilized peroxidase, which can act as an outer interference barrier by preoxidation of the unwanted species before reaching the underlying polarized transducer. Brahim et al. [154] demonstrated a very effective interference screening membrane for amperometric biosensors by combining extensively oxidized polypyrrole within a cross-linked p(HEMA)-based hydrogel. Even though the polypyrrole component was devoid of electroactivity, the resulting polymer composite increased the selectivity of the hydrogel membrane by a combination of size exclusion and anion exclusion imparted by the conditioned conetwork polymer. Perhaps the obvious advantage of using polymer multilayers in biosensor design is that they offer a solution for enzyme immobilization when it is otherwise not feasible in a single-polymer layer, as is the case with polyaniline biosensors whose electropolymerization conditions in acid media tend to significantly compromise the activity of immobilized enzymes [155].
12.3.5 Immobilization of Redox Mediators Three general techniques have been pursued to incorporate redox mediators into conducting polymers for the generation of more efficient biosensors: (i) entrapment of the mediator during polymer electrosynthesis, (ii) covalent linking of the mediator to the enzyme followed by incorporation of the derivatized enzyme during polymer electrosynthesis, and (iii) covalent linking of the mediator either to the monomer followed by electrosynthesis of the derivatized conducting polymer, or covalent linking of the mediator to the preformed conducting polymer. The first immobilization technique is the most straightforward; the mediator species is dissolved in a solution containing the monomer such that the mediator exists predominantly in its negatively charged
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state. During the electropolymerization process, the conducting polymer is grown and simultaneously incorporates the mediator into the network as a counteranion. The entrapped mediator species should remain within the conducting polymer network yet be sufficiently mobile to function as an electron shuttle between the enzyme and electrode surface. Examples of mediators immobilized within conducting polymers using this technique include ferrocyanide [156], ferrocene derivatives [157], phenazine [158], and quinones [159]. One problem revealed by this approach is the potential for egress of the counteranion mediator following redox cycling of the conducting polymer. To overcome excessive mediator loss through leeching out of the conducting polymer matrix, an alternative technique is to covalently link the mediator to the enzyme molecule to be entrapped. Apart from preventing mediator loss, this procedure improves the efficiency of the chargetransfer process by immobilizing several mediator molecules to each enzyme molecule. Moreover, this increases the probability of interaction between these two species since the local distance is now reduced [160]. Direct covalent linking of the mediator to the conducting polymer is a less attempted technique than (i) and (ii) above. This is because of the likely compromise of the electron-shuttling capacity of the confined mediator species to the active site of the entrapped enzyme and the transducer surface. Nevertheless, this strategy has produced amperometric biosensors for glucose and cytochrome c [161,162]. The glucose biosensor was fabricated by covalently linking osmium complexes to the polypyrrole backbone via flexible spacer chains and then entrapping PQQ-dependent glucose dehydrogenase within the matrix. The proposed underlying mechanism was that the electron transfer (ET) distance between the enzyme’s active site and the electrode surface, now modified with this mediator-containing polymer, divided into a layered sequence of several self-exchange ET reactions that allowed ‘‘electron hopping’’ to occur. The amperometric detection of cytochrome c in the latter biosensor was achieved by the use of ferrocene-derivatized polythiophene. Both thiophene monomer and terthiophene were covalently linked with ferrocene mediator to produce novel homopolymer and copolymer materials. Both materials revealed cyclic voltammograms with redox peaks that were assigned to the mediator and the conducting polymer backbone constituents, respectively. The ability of these materials for direct electrical communication with proteins was demonstrated with concomitant cytochrome c oxidation and ferrocene reduction, which was not feasible with the underivatized conducting polymer: Fc $ Fþ c þe þ Cyt c þ Fþ c ! Fc þ Cyt c
12.3.6 Derivatization of Monomers before Conducting Polymer Electrosynthesis Examination of the chemical structure and properties of the pyrrole monomer reveals that the nitrogen atom is amenable to chemical substitution reactions [163]. Apart from these N-substitutions, derivatizations at positions 3 or 4 on the pyrrole ring can be performed. Depending on the site of derivatization, the resulting derivatized monomers may require higher polymerization charge density that results in a polymer film with reduced conductivity (N-substituted), or conversely, may undergo electropolymerization at lower anodic potentials and lead to the formation of conducting polymer matrices with enhanced electronic conductivity (substituted third and fourth heteroatoms). A novel group of electroactive polymers, the polyazines, was demonstrated by Schmidt and coworkers [164] (Figure 12.11). This group includes the widely known indicator molecules methylene blue, thionine, and Meldola blue, which have been previously exploited for their mediating properties as well as an immobilizing matrix in the fabrication of biosensors for glucose, choline, and hydrogen peroxide [165].
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Cosnier et al. [166] has developed electropolymerizable materials of a dicarbazole-derivative functionalized by N R3 N-hydroxysuccinimide and pentafluorophenoxy groups [166]. The subsequent chemical functionalization of the poly(dicarbazole) film was easily performed by successive R2 X R1 immersions in aqueous enzyme and mediator solutions. FIGURE 12.11 Structure of the repeatThese derivatized, bioactive conducting polymer films were ing unit of polyazine. (Reprinted from demonstrated as sensing layers for catechol. Khan, G.F., Kobatake, E., Ikariyama, Y., A number of other conducting polymer derivatives have and Aizawa, M., Anal. Chim. Acta, 281, been used for immobilizing enzymes and redox mediators 527, 1993. With permission.) for biosensor applications. Among these polymers are poly (tyramine) [167] and poly(1-(5-aminonaphthylethanoic acid) [168], which have accessible amine and carboxyl functionalities, respectively, for covalently linking to complementary groups on available amino acid residues of enzymes. Polypyrrole and polythiophene derivatives have also been widely exploited for covalent immobilization of enzymes, which then function as the sensing layers in efficient biosensors [169,170]. Another strategy that has attempted to improve the use of conducting polymer–based matrices for biomolecule immobilization and consequent stabilization of bioactivity is the fabrication of amphiphilic monomers of pyrrole [171,172]. These derivatives are synthesized to contain polar and nonpolar zones that result in stable cationic polymer films. This approach of enzyme immobilization is first based on the solubilization of biomolecule in an aqueous dispersion of the amphiphilic pyrrole monomer. Subsequently, the aqueous mixture is spread and dried on an electrode surface. The electropolymerization of the adsorbed monomers in an aqueous electrolyte provides the irreversible entrapment of biomolecule in the resulting polypyrrolic matrix. The main advantage of this procedure of enzyme immobilization is the possibility to control and improve the composition of the biomolecule–polymer layer, particularly with respect to biomolecule loading. Further improvements to this strategy have been to include hydrophilic clay particles within the conducting amphiphilic polymer, which has been demonstrated to preserve enzyme bioactivity [173], and encapsulation of enzymes in liposomes before immobilizing to the amphiphilic monomer [174], which typically yields improvements to biosensor sensitivity. H
12.3.7 Direct Electrical Modulation of Enzymes To exert specific electrochemical control and modulation over the catalyzing actions of enzymes has been a long sought after goal in bioelectrochemistry and biotechnology. Conducting polymers have been demonstrated to affect, or switch, some degree of regulation of enzyme activity in the following ways: (i) switching of the regeneration of the components in the electron-transfer chain between the active site of the enzyme and the transducer surface, particularly so with the inclusion of coenzymes and mediators [175], (ii) direct switching of the enzyme activity across the redox state of the conducting polymer through direct electrical communication between enzyme and polymer [176], and (iii) switching of the availability of critical species in the enzymatic reaction when ionic species such as substrates or coenzymes are involved [177]. One investigative study that has raised considerable queries is the report of direct electrical communication between the active site of the enzyme and the conducting polymer when the enzyme is immobilized in polypyrrole microtubules. These microtubules were produced by electropolymerization of the pyrrole monomer inside the pores of a microporous filtration membrane. This configuration is reported to favor direct electron transfer across the polymer structure as well as direct reoxidation of the enzyme at lower potentials than typically used, hence promoting increased selectivity of the resulting amperometric biosensor [178]. Another report investigating the same system claims that it is the underlying platinum metal rather than polypyrrole tubules that is responsible for the observed direct enzyme reoxidation [179].
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12.3.8 Electroactive Polymer Hydrogels The first reported investigation of the synthesis of composite materials comprising conducting polymers and hydrogels was in 1997 by Wallace and coworkers [180]. The primary objective behind this study was to fabricate polymer composites that would result in materials with enhanced porosities and iontransport properties to be used for controlled delivery applications through electrochemical-stimulated release of analytes. Polypyrrole and polyaniline were galvanostatically electrosynthesized with varying incorporated counterions within polyacrylamide hydrogels in a custom-built cylindrical gel cell (Figure 12.12A). The resulting polymer composites possessed very high water contents (ca. 90%) and could be dehydrated and subsequently rehydrated back to 80% of the initial water content. This characteristic was suggestive of an open porous structure associated with the composite hydrogel network (Figure 12.12B). Subsequent electrostimulated-controlled release of large molecules (calcon) that were previously doped into the conducting polymer also supported an open, porous network structure, with higher released quantities and rates of release observed for the hydrogel-conducting polymer composites compared to the pristine polymer. The conducting polymer component in these composites retained its redox electroactive behavior when electropolymerized at relatively low polymerization times. The same group also demonstrated blending of conducting polymer colloids with processable hydrogels of poly(acrylic acid). Both soluble and extrudable polymer composites were fabricated. In both systems, the conducting polymer component displayed electroactive behavior with high equilibrium water contents. However, the scientists did not demonstrate the use of such hydrogel-conducting polymeric composites for biosensor applications. Another class of hydrogel-conducting polymer composite material, poly(3,4-ethylenedioxythiophene)–poly(styrenesulfonate) (PEDOT–PSS) in combination with the redox hydrogel poly(4-vinylpyridine) with coordinated osmium, was synthesized primarily for use as a matrix for bioimmobilization
Reference electrode
Hydrogel network
Salt bridge
Electrochemical cell
Monomer solution
Hydrogel conducting polymer network
RVC auxiliary electrode Platinum mesh working electrode 10 mm 25 mm
Polyacrylamide hydrogel
40 mm
(A)
(B)
FIGURE 12.12 (A) Concentric electrode configuration gel cell used by Wallace and coworkers for electrosynthesis of conducting polymer within hydrogels and (B) Schematic representation of the open porous structure of the polymer gel that is preserved after growth of the conducting polymer component. (From Small, C.J., C.O. Too, and G.G. Wallace, Polym. Gels. Networks, 5, 251, 1997.)
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for the fabrication of enzyme electrodes [181]. The construction of a hydrogen peroxide electrode by immobilizing HRP within the polymer composite was demonstrated. The electrode was capable of amperometrically detecting hydrogen peroxide at substantially lower polarizing potentials (þ0.025 V versus Ag=AgCl, 3 M Cl) than typically used. Of particular interest was the highly porous nature of the resulting hydrogel composite, a feature that may render the material very applicable in the fabrication of nerve electrodes and cellular interfaces wherein such an interface closer mimics the biological structures thus, potentially enabling cells to grow even inside the electrode structure. There have been several studies conducted on the fabrication of polymer blends of hydrogels such as poly(methyl methacrylate) (PMMA), poly(vinyl methyl ether) (PVME), and p(HEMA) with electronically conducting polymers such as polypyrrole, polyaniline, or their derivatives primarily for the construction of artificial muscle fibers [182,183]. With respect to biosensing applications, the group of Guiseppi-Elie has demonstrated the synthesis of ‘‘biosmart’’ hydrogel composite materials, consisting of cross-linked p(HEMA) with incorporated polypyrrole and polyaniline chains that are tailored biospecific by immobilization of enzyme molecules. Brahim et al. [184] fabricated biosmart polypyrrole–p(HEMA) composites to function as sensing membranes for clinically important amperometric biosensors. A monomer cocktail containing, among other components, the methacrylate monomers, pyrrole or aniline monomer, and photoinitiator was first irradiated by UV light to effect polymerization of the hydrogel components. This was immediately followed by potentiostatic electropolymerization of the pyrrole–aniline monomer in a PBKCl solution saturated with further monomer. The resulting polymer optical transition from transparent to dark green opaqueness visually confirmed growth of the conducting polymer component within the cross-linked hydrogel matrix (Figure 12.13). Amperometric enzyme biosensors for the detection of glucose, cholesterol, and galactose were demonstrated, each possessing extensive linear dynamic response ranges, high sensitivities, and prolonged storage stabilities. The included and extensively oxidized conducting polymer component within these crosslinked hydrogel composites was shown to effectively screen against common electrooxidizable interfering species, both anionic and neutral. Further electrochemical investigations conducted on these p(HEMA)–PPy films revealed that the incorporated conducting polymer was actually devoid of electroactivity [154] after the extensive polarization treatment at þ0.7 V (versus Ag=AgCl, 3 M Cl). It was proposed that the continuous application of this oxidizing potential while the film was immersed in a
FIGURE 12.13 Change in opaqueness of electrode-supported hydrogel (transparent gel without CEP) upon growth of the conducting polymer component (dark green gel with CEP) within the gel network.
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solution of phosphate buffer could result in the disruption of conjugation throughout the polymer by the generation of anionic groups with high electron density on the polymer backbone. A novel polymer composite material consisting of a water-dispersed complex of polypyrrole doped with polystyrenesulfonate and embedded in polyacrylamide hydrogel was prepared and evaluated as a matrix for enzyme immobilization [185]. The enzyme glucose oxidase was entrapped in the polymer complex by inclusion in the aqueous phase of the emulsion polymerization synthesis process. The resulting bioactive microparticles, having average diameter between 3.5 and 7.0 mm, were placed on the surface of cleaned platinum electrodes and compressed using dialysis membrane. The polymer-modified electrodes were then used as amperometric biosensors for the determination of glucose under both aerated and deaerated conditions with rapid response times and efficient screening of interferents.
12.4
Biotransducer Devices and Biosensor Systems
12.4.1 Microfabricated Array Electrodes Microelectrodes, as the name suggests, define a category of electrodes and devices having active working features with dimensions of micrometers (106 m) or less (nanometers, 109 m) and are produced by the processes of micro- and nanolithography. Microfabricated array electrodes have become the mainstay of CEP biotransducer research and development, in part because of their commercial availability from ABTECH Scientific, Inc. Amongst these devices, shown in Figure 12.14, are the interdigitated microsensor electrodes (IMEs), independently addressable microband electrodes (IAMEs), independently addressable interdigitated microsensor electrodes (IAIMEs), microdisc electrode arrays (MDEAs), independently addressable microdisc electrode arrays (IAMDEAs), and inverted or buried gate field-
Substrate
Conductor
(A) Region of Si3N4 Substrate Conductor interdigitation Layer
Microbands
Si3N4 Layer
(B)
1 cm
Bonding pads Reference electrode
Microdisc array
MDEA 5037 Counter electrode
2 cm
Silicon nitride
Bonding pad (D)
Borosilicate glass substrate ABTECH MDEA 5037
Gold or platinum
Si3N4 pasivation layer
(E)
(C)
FIGURE 12.14 Schematic of various microfabricated electrode array devices manufactured by ABTECH Scientific, Inc. (A) Interdigitated microsensor electrodes (IMEs), (B) independently addressable microband electrodes (IAMEs), (C) independently addressable interdigitated microsensor electrodes (IAIMEs), (D) microdisc electrode array (MDEA), and (E) electrochemical cell-on-a-chip MDEA (ECC MDEA 5037).
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effect devices (FETs). These devices offer advantages over conventional larger working electrodes within biosensors since they experience hemispherical (radial) solute diffusional profiles as opposed to linear diffusional profiles (Figure 12.15) is this phenomenon that can impart stir (convectional mass transport) independence to sensor responses, while also offering lowered limits of detection and faster attainment of equilibrium reactions, resulting in more rapid response times. Also, microelectrodes can often be used in high-resistance media due to the low-operational currents typically encountered. Although it may be argued that individual microelectrodes offer very small responses, one approach for overcoming this problem is to use many microelectrodes together in the form of an array to allow a cumulative and so larger response to be measured. Electropolymerized-conducting polymers have been extensively investigated as materials for solid state and electrochemical biosensor devices possessing either of these two designs above. The interdigitated array (IDA) electrode coated with an electroconductive polymer has been demonstrated to serve as a conductimetric, pH-sensitive, sensor device. Nishizawa and Uchida [186] fabricated a penicillin sensor by first potentiostatically depositing a thin film of polypyrrole over the digits of an IDA that were pretreated with octadecyltriethoxysilane to create hydrophobic areas on the surrounding nonconducting, glass substrate. The polypyrrole-coated electrode was then covered with a top layer of penicillinase immobilized in BSA–glutaraldehyde mix. Penicillinase catalyzes the hydrolysis of penicillin to penicilloic acid and acidifies the PPy ultrathin film. The resulting ohmic current was seen to increase linearly with concentration of penicillin up to 7 mM. Since many enzymes bring about pH changes through their catalytic reactions, this demonstrated principle would be widely applicable to the fabrication of enzymebased microelectrochemical devices that detect biologically important molecules. The oxidation or reduction of conducting polymers via enzyme reaction is another possible principle of the impedimetric biosensor. It was reported previously that the microarray electrode coated with the copolymer of pyrrole and N-methylpyrrole containing diaphorase worked as a switching device, which showed ‘‘on–off ’’ response to the presence of NADH [187]. Bartlett and coworkers have expanded this
(A) Microdisc
Microelectrodes
(B)
Planar
Macroelectrodes
FIGURE 12.15 Lines of flux that predominate at different electrode dimensions (A) radial diffusion at microelectrodes and (B) linear diffusion at macroelectrodes.
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idea to fabrication of sophisticated bioelectrochemical devices responsive to glucose [188] and NADH [189]. In the former case, glucose oxidase was immobilized in an electropolymerized film of insulating poly(l,2-diaminobenzene) grown on top of a polyaniline base film that was electropolymerized over a pair of screen-printed carbon microband electrodes (Figure 12.16) [190]. These were fabricated on a PVC base by screen printing successive layers of carbon followed by two layers of dielectric, by another layer of carbon, and finally by two layers of dielectric. The dielectric provided insulation between the two carbon microband electrodes, which were typically 0.5 cm long and 10–15 mm wide with the gap 20–25 mm wide. Polyaniline was chosen as the active conducting polymer material because, unlike poly (pyrrole) or poly(N-methylpyrrole), it is stable in the presence of hydrogen peroxide and because it can be switched from an oxidized, insulating state to a less oxidized, conducting state. This corresponds to the change from the pernigraniline form to the emeralidine form of the material. While there are currently several fabrication approaches to produce microelectrode arrays, techniques such as photolithography or laser ablation have to date proved cost-prohibitive for the mass production of disposable microsensor strips for commercial applications. A novel, patented procedure created by the research group of Higson [191] at the Institute of Bioscience and Technology, Cranefield University at Silsoe, UK, allows the fabrication of densely populated (up to 2 105 microelectrode elements per unit area) microelectrode arrays for biosensor applications via a simple and inexpensive sonochemical ablation approach. In this technique, a thin (30–40 nm thick), insulating layer of poly-
Device OFF A GOx
i=0
Poly(1,2−diaminobenzene) Potentiostat at +0.5 V versus SCE
Glucose
Poly(aniline)
Carbon electrodes (A) Device ON
(B)
A i >0
FIGURE 12.16 (A) Enzyme switch responsive to glucose. The enzyme glucose oxidase is immobilized in a thininsulating film of poly(l,2-diaminobenzene) deposited on top of the poly(aniline) film and (B) operation of the switch. Upon exposure to glucose, the switch is turned from ‘‘off ’’ to ‘‘on.’’ The switch is reset by holding the film at þ0.5 V versus SCE in a background buffer solution. (Reprinted from Nishizawa, M. and Uchida, I., Electrochim. Acta, 44, 3629, 1999. With permission.)
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Hemispherical diffusion
Insulating polymer
Array
Underlying substrate (A)
(B)
FIGURE 12.17 Deposition of polydiaminobenzene to form an insulating film. (A) Schematic of polymer-insulated electrode and (B) schematic of microelectrode array. (Reprinted from Matsue, T., Nishizawa, M., Sawaguchi, T., and Uchida, T., J. Chem. Soc., Chem. Commun., 15, 1029, 1991. with permission)
diaminobenzene was first electropolymerized (from 1,2 diaminobenzene solution) over a planar gold electrode surface. Sonochemically fabricated microelectrodes were subsequently prepared via the ablation of polydiaminobenzene film using an ultrasonic frequency of 25 kHz. Ultrasound in this kHz range passing through a solvent can lead to the formation of microjets of solvent that travel at velocities up to several hundred meters per second, causing ablation and cavitation of the soft, solid polymer surface (Figure 12.17). These cavities were observed by an SEM to be polydisperse with an average diameter of 3 mm or less. To demonstrate the feasibility of these sonochemically fabricated microelectrode arrays for biosensors, Higson and coworkers [191] electrochemically coated the micropores with polyaniline containing glucose oxidase, grown potentiodynamically from a phthalate buffer solution containing the monomer and enzyme. Such bioactive microelectrode arrays were demonstrated for amperometric and impedimetric transduction of glucose, with consequent enhanced sensitivities and limits of analyte detection. It was also stated that further enzymatic biosensor systems employing the novel microelectrode arrays were developed for the detection and quantification of ethanol, oxalate, several pesticides, as well as for the detection of DNA hybridization and the fabrication of affinity antibody–antigen biosensors.
12.4.2 Electroactive Polymer Sensor Interrogation System for Conductimetric Response and Impedimetric Response EPSIS represents a unique analytical tool designed for research and development of chemical and biological sensor devices, instruments, and systems based on electroconductive polymer sensor technology. Electroconductive polymers form effective conductimetric transducers for chemical and biological sensors. EPSIS combines potentiometric, potentiostatic, and superior pulsed DC chronocoulometric capabilities into a powerful and versatile analytical detection and measurement method that is unique to the determination of conductimetric chemical and biological sensor responses of electroconductive polymers. The EPSIS analytical method involves the following three distinct phases: 1. Preinitialization—An Undisturbed Open Circuit Potential Measurement. This step measures, stores, and presents to the screen the interfacial potential of the electroconductive transducer in its electrolyte or test environment. This measurement is made of the complete polymer film relative to a reversible Ag=AgCl, Cl electrode that also contacts the electrolyte. The unperturbed open circuit potential is diagnostic of the redox state of the device and conveys information about the integrity of the transducer. 2. Initialization—A Conditioning Electrolysis. EPSIS applies a user-specified potential to the electroconductive polymer transducer for a user-specified duration or a user-specified limiting current. Electroconductive polymer transducers possess redox active sites that are present in varying amounts of oxidized and reduced forms. By applying an initialization potential, EPSIS
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Resistance
109 Ohms 10−2
Current
+0.5
mA −0.5 +1.0
Poise potential
v +1.0 Preinitialization
Time base Initialization
Interrogation
FIGURE 12.18 Schematic illustration of the EPSIS interrogation phases indicating resistance, current, and potential responses. (Reprinted from Guiseppi-Elie, A., Wallace, G.G., and Malsue, T., Handbook of Conducting Polymers, Marcel Dekker, New York, 1998. With permission.)
fixes the initial redox composition of the transducer, thereby also fixing its initial electrical conductivity and sensitivity toward biologically induced redox changes. 3. Interrogation—A Nonpertubating Measurement of the Time Dependence of the Electrical Conductivity of the Transducer. Interrogation is itself broken into a sequence of events designed to reveal the time dependence of the electrical conductivity of the transducer as it responds to biologically induced redox changes. These four events include the application of a nonpertubating but interrogating voltage pulse of 10–50 mV applied between the fingers of the IME device; measurement of subsequent transducer conductivity by integration of its current such that the measured charge is directly proportional to the transducer’s conductivity during the pulse period; a float period whereby the potential is completely withdrawn from the fingers of the device and the device is allowed to spontaneously respond to biologically induced redox changes; and remeasurement of the open circuit potential, which then becomes the basis for the application of the subsequent voltage pulse. The three phases of preinitialization, initialization, and interrogation are illustrated schematically in Figure 12.18. The four sensor interrogation steps of pulse application, conductivity measurement, float period, and open circuit potential measurement are repeated for a user-defined number of cycles to produce a sensor response curve. Each set of four such steps produces a single-datum point of conductivity and open circuit potential data. Several such points obtained over a period of time produces a response curve. The response curve captures the change in conductivity of the transducer as a function of time following initialization. The group of Guiseppi-Elie [135] demonstrated the design, fabrication, and operation of polypyrrolebased conductimetric biosensors for glucose and urea using the EPSIS methodology. The principle of operation of the glucose biosensor was based on two phenomena: (i) the very strong binding (affinity) between biotin and streptavidin, which covalently immobilizes the glucose oxidase enzyme to the PPy– IME substrate and confers glucose sensitivity, and (ii) the oxidation of polypyrrole by enzymatically generated H2O2, shown coupled in the following equations:
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GOx
b-D-Glucose þ O2 ! Gluconolactone þ H2 O2 H2 O2 þ 2Hþ Cl þ PPy ! 2H2 O þ (PPy þþ 2Cl ) The H2O2 causes oxidation of PPy, resulting in a concomitant and stoichiometric change in inherent conductivity, which was recorded by the EPSIS instrument. The response of the conductimetric glucose biosensor to different glucose concentrations was effected by a corresponding increase in conductivity of the PPy membrane over all glucose concentrations investigated. The rate of conductivity change was seen to increase with glucose concentration from approximately 4 106 S=cm=s at 100 mM glucose to nearly 9 106 S=cm=s in response to 600 mM glucose. The EPSIS was also used to demonstrate a conductimetric urea biosensor employing the biotinylated PPy membrane grown over the IMEs and incorporating the enzyme urease through a biotin–urease conjugate involving streptavidin coupling. In this system the change in conductivity of PPy, measured in deionized (DI) water, is effected through the following enzymatic reaction: Urease
þ Urea þ 3H2 O ! HCO 3 þ 2NH4 þ OH
The enzymatic products, ammonia and OH ions, serve to compensate the formation of charge carriers within the polymer, resulting in an ability to decrease the apparent polymer conductivity. The conductimetric response pattern was, however, complicated by a second opposing phenomenon; the increasing concentration of anions at the polymer electrolyte interface that serve to increase apparent conductivity via the availability of doping anions. With this configuration therefore, the polypyrrole device displayed a ‘‘step’’ in its response that corresponds to an ON from an OFF condition at a particular urea concentration. The authors anticipated that such a threshold response may be engineered by changing the activity of the enzyme within the polymer membrane. Such a detector may be used to monitor the urea content of the dialysate from patients under urea dialysis.
12.4.3 Microcantilevers Piezoelectric materials incorporating conducting polymers have been investigated for use in various applications, such as sensors and actuators, due to their flexibility, ease of processing, lightweight character, and low cost. Joo and coworkers [192] recently fabricated a bimorph cantilever incorporating the piezoelectric polymer (poly(vinylidenefluoride), PVDF, b-phase) for the active layers and the highly conducting polymer (PEDOT–PSS) treated with a dimethyl sulfoxide (DMSO) solvent for the electrodes (Figure 12.19). The PVDF films were modified so as to have high adhesion at the interface between the PVDF and the PEDOT=PSS (DMSO) film by using an ion-assisted-reaction (IAR) method. The tip displacement of the cantilevers was measured at a resonance frequency of 27 Hz, and the deformation of the PVDF film with the IAR-treated PEDOT=PSS (DMSO) electrodes was measured to be as much as 7 mm when under the stimulus of an input voltage of 40 Vrms. This displacement was found to be higher than that with PEDOT=PSS or pristine inorganic electrodes at the same interrogating frequency and voltage. The cantilevers made with indium tin oxide (ITO) or platinum (Pt) electrodes became damaged after operating the devices at a high frequency or a high input power. The PVDF cantilevers made with the PEDOT=PSS (DMSO) electrodes, however, were observed to be electrically and mechanically durable when operating at both high input voltage and high operational frequency.
12.4.4 Biocompatiblity for In Vivo Sensing Biocompatibility in its broadest sense refers to the ability of a material to perform with an appropriate host response in a specific situation. Any foreign material once implanted within an organism triggers a cascade of reactions called the immune or host response, which are part of the organism’s defense mechanism. Several studies have been performed to evaluate the biocompatibility of polymers [193],
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Bas
elin
PVDF Electrode Epoxy
Dipole
(A)
e
Tip displacement
de
No
(B)
FIGURE 12.19 (A) Schematic diagram of the parallel-type bimorph cantilever and (B) Photograph of the actuation of PEDOT=PSS polymer–based cantilevers. (From Lee, C.S., J. Joo, S. Han, and S.K. Koh, Sens. Actuators A, 373, 2005.)
including the use of conducting polymers for in vivo applications. Polypyrrole is an interesting candidate for tissue-engineering applications, by virtue of its inherent electrical conductivity, the ease with which one can control crucial surface properties such as wettability and charge density, and its compatibility with mammalian cells. In addition, polypyrrole is quite unique in the class of biomaterials because it is tissue compatible and exhibits longevity when implanted, has been demonstrated to improve the regeneration of tissues, can serve as suitable matrices for the growth and proliferation of cell cultures, and can be doped with subsequent releasable biomolecules [194]. In one study by Langer and coworkers [195], the utility of the electrically conductive polymer, polypyrrole, as a substrate to enhance nerve cell differentiation in culture was evaluated with the hope of ultimately using electrically conducting polymers to stimulate in vivo nerve regeneration. The researchers convincingly demonstrated that polypyrrole is a suitable material for in vitro nerve cell culture (rat PC 12 cells) and that application of an electric stimulus through the conducting polymer enhances neurite outgrowth. In addition, they showed that polypyrrole does not elicit an adverse tissue response when implanted in both rat subcutaneous and muscular tissue (Figure 12.20). The biocompatibility of polyaniline films in the emeraldine, nigraniline, and leucoemeraldine intrinsic oxidation states were assessed through subcutaneous implantation of 1.0 1.5 cm slabs into Sprague Dawley rats for periods ranging from 19 to 90 weeks. Upon explantation of the PANI films, there was no observed inflammation or infection in the surrounding tissue over the entire implantation period. These and associated parallel polymer surface characterization studies clearly demonstrated that these conducting polymers are sufficiently biocompatible, at least when implanted in the dorsal region of the skin, to be used for biomedical applications. In another study designed to investigate the biocompatibility of in situ utilization of biomaterials [193], nonresorbable materials separately composed of polypyrrole, polyaniline, and polyimide were fabricated along with resorbable materials (PLLA-PDXO-PLLA) for tissue-engineering applications and to investigate their overall tissue tolerance and cellular interactions. These nonresorbable, intrinsically conductive polymers were conceived for biosensor applications and implantable drug delivery systems. It was observed that the cells in contact with the resorbable material appeared to be capable of migratory-regenerative aspects in vitro and also exhibited good compatibility in vivo, whereas the nonresorbable materials, which are designed to remain in situ in vivo, were seen to have the potential to represent an adverse factor (inflammation, fibrotic reactions) that correlated with some aspects of cell behavior in vitro.
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A
C
PP
PLGA B
FIGURE 12.20 PC-12 cell differentiation on polypyrrole (A) without and (B) with application of an electric potential. PC-12 cells were grown on PPy for 24 h in the presence of NGF, then exposed to electrical stimulation (100 mV) across the polymer film, S (B). Images were acquired 24 h after stimulation. Cells grown for 48 h but not subjected to electrical stimulation, NS, are shown for comparison (A). Bar ¼ 100 mm. (C) Histology of tissue response to polypyrrole. Shown is a histological tissue section stained with hemotoxylin and eosin of a PPy=PLGA disk that had been implanted in rat muscular tissue for 2 weeks. The PPy is seen as a black film (left), and the PLGA is almost transparent (right). The PPy and PLGA films separated during histological processing, explaining the gap in the center of the image. (From Schmidt, C.E., V.R. Shastri, J.P. Vacanti, and R. Langer, Proc. Natl. Acad. Sci. USA, 94, 8948, 1997.)
While several conducting polymers have been demonstrated to be biocompatible, there is still the concern of the longevity of these materials and their potential long-term impact within the in vivo environment such as inducing chronic inflammation, since such polymers are nondegradable. Schmidt et al. [196] have begun to tackle this potential concern by synthesizing novel biodegradable electrically conducting polymers (BECPs) composed of monomeric sequences of alternating units of pyrrole and thiophene moieties capped on either side by degradable ester linkages and aliphatic linkages (Figure 12.21). Human neuroblastoma cells cultured in vitro on BECP films demonstrated attachment and neurite extension after 1 d and significant proliferation after 8 d, indicating good cell compatibility. In vivo biocompatibility was characterized by subcutaneous implantation into rats for 14 and 29 d, with only mild inflammation observed and significant tissue infiltration beyond the tissue–polymer interface. In another study, Inoue et al. [197] reported a polypyrrole derivative–based glucose sensor fabricated by the electropolymerization of 1-(6-D-gluconamidohexyl)pyrrole (GHP) and 6-(1-pyrrolyl)hexylphosphatidylcholine (PPC) on a platinum wire electrode precoated with polydimethylsiloxane. While the response of the electrode prepared from PPC was satisfactory, no response was observed from the electrode of GHP. By the addition of Nafion into the precoating solution however, the improvement of sensor sensitivity occurred and the electrode prepared from GHP also showed clear sensor response. Moreover, the introduction of Nafion was effective to improve the long-term stability of the enzymeimmobilized electrodes. They also reported that the electropolymerized membrane prevented plasma from getting absorbed by the sensing membrane when the electrodes were used to measure glucose in blood samples. While various BioMEMs (Bio-microelectromechanical systems) applications have been demonstrated using conductive polymers, particularly polypyrrole, there has yet been no reported use of conducting polymers solely for implantable biosensor devices. However, recent material advances such as those
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Electroactive oligomeric unit O O
H N
S
H N
O
O
O O
O
O
4
n
Degradable ester linkages
Aliphatic linker
FIGURE 12.21 Schematic of biodegradable electrically conducting polymer (BECP) illustrating its key components: a conducting pyrrole–thiophene–pyrrole oligomer, degradable ester linkages, and an aliphatic linker. (From Rivers, T.J., T.W. Hudson, and C.E. Schmidt, Adv. Funct. Mater., 12, 33, 2002.)
described above suggest that the future for conducting polymers as sensing materials for in vivo biosensors seems promising and certain.
12.4.5 Electronic Noses Since the pioneering work of Persaud and Dodd [198] in the mid-1980s with conducting polypyrrole, the development and application of conducting polymer–based sensor arrays, or electronic noses, have increased exponentially into diverse areas such as odor sensing for gases [199,200], recognition of spoilage bacteria and yeasts in milk [201], water quality control [202], identification of fruit cultivar [203], and characterization of wines [204] and olive oil [205]. One quite intriguing application was the use of polyaniline-based polymers integrated into an electronic nose platform for discriminating predialysis from postdialysis blood as well as control blood for the management of renal failure [206]. Additionally, the manufacture of commercial electronic nose sensor devices have been placed with advanced microfabrication technologies such as photolithography, ink-jet printing, and microcontact printing. Novel strategies have been pursued in attempts to preserve the stability of the incorporated polymer components for both CEPs as well as insulating polymers. One such approach is the inclusion of carbon black particles into electrically insulating thin-film polymers [207]. The attractiveness with this tactic is that the conductive element is very stable carbon black, and the chemical diversity required for the sensor array can be readily obtained by using different organic polymers as the insulating phase for the carbon black polymer composites. In a recent study employing a 32-sensor array containing this base polymer system incorporated into a handheld unit, researchers demonstrated good correlation for expired breath components determined by the electronic nose sensor and the same exhaled volatiles analyzed for using a clinical pneumonia test [208]. In another study, it was also found that this type of sensor array could be used to resolve common organic solvents, including molecules of different classes (e.g., aromatics against alcohols) as well as those within a particular class (e.g., benzene versus toluene or methanol versus ethanol).
12.4.6 Nanowire Arrays In a dimensional realm coined by the ‘‘nano’’ buzzword and populated with nanomaterials dominated by semiconductor and metal nanowires (NWs) and CNTs, conducting polymers are emerging as a promising material for the synthesis of nanostructured materials and devices. They are particularly appealing because they exhibit electrical, electronic, magnetic, and optical properties similar to metals or semiconductors while retaining their flexibility, ease of processing, and modifiable electrical conductivity. Their porous structures are very amenable to entrapping biomolecules. A variety of conducting polymers have shown promise as sensor materials, including biosensors, because their properties can
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Anode (−)
PANI Nanowire
Anode (−)
Cathode (+)
Anode (−)
2 µM
Acc V Spot Magn Det WD 10.0 kV 3.0 30000x SE 21.2
Cathode (+)
PPY Nanowires gn xx
(A)
Cathode (+)
Det SE
10 µM
WD 16.5
(B)
FIGURE 12.22 (A) SEM image of a 100 nm wide by 4 mm long polyaniline (PANI) nanowire. Scale bar: 2 mm and (B) SEM image of two 200 nm wide by 2.5 mm long PPy nanowires separated by 10 mm deposited one at a time. Scale bar: 10 mm. (From Ramanathan, K., M.A., Bangar, M. Yun, W. Chen, A. Mulchandani, and N.V. Myung, Nanoletters, 4, 1237, 2004.)
be tailored to detect a wide range of chemical compounds. Ironically, the majority of limitations associated with processing of NWs and CNTs for development of high-density sensor arrays also plague the use of conducting polymers for these applications. Recently, Ramanathan et al. [209] described a facile technique for the fabrication of individually addressable, conducting polypyrrole and polyaniline nanowire arrays of controlled dimension and high aspect ratio (100 nm wide by up to 13 mm long) created in channels on insulating surfaces (Figure 12.22). This feat was achieved via simple threeelectrode electrodeposition (galvanostatically) of polymer from starting monomer and dopant solution placed between electrodes. In addition, the ability to create ‘‘arrays’’ of conducting polymer nanowires of same or different materials on the same chip was demonstrated. The feasibility of such conducting polymer nano-assemblies to function as sensors was demonstrated by measuring inherent changes to the conductivity of the polymer nanowires produced in response to varying pH stimuli. Moreover, the benign-operating conditions for electropolymerization make this process ideal for the fabrication of nanobiosensors by direct deposition of conducting polymer nanowires with embedded bioreceptors in one step rather than the multiple steps presently required in surface-modified nanowires and CNTs.
References 1. Castillo, J., S. Ga´spa´r, S. Leth, M. Niculescu, A. Mortari, I. Bontidean,V. Soukharev, S.A. Dorneanu, A.D. Ryabov, and E. Cso¨regi. 2004. Biosensors for life quality: Design, development and applications. Sens Actuators B 102:179. 2. Hofstadler, S.A., R. Sampath, L.B. Blyn, M.W. Eshoo, T.A. Hall, Y. Jiang, J.J. Drader, J.C. Hannis, K.A. Sannes-Lowery, L.L. Cummins, B. Libby, D.J. Walcott, A. Schink, C. Massire, R. Ranken, J. Gutierrez, S. Manalili, C. Ivy, R. Melton, H. Levene, G. Barrett-Wilt, F. Li, V. Zapp, N. White, V. Samant, J.A. McNeil, D. Knize, D. Robbins, K. Rudnick, A. Desai, E. Moradi, and D.J. Ecker. 2005. TIGER: The universal biosensor. Int J Mass Spectrom 242:23. 3. (a) Shirakawa, H., E.J. Louis, A.G. MacDiarmid, C.K. Chiang, and A.J. Heeger. 1977. Synthesis of electrically conducting organic polymers: Halogen derivatives of polyacetylene, (CH)x. J Chem Soc Chem Commun 578; (b) Chiang, C.K., C.R. Fincher Jr., Y.W. Park, A.J. Heeger, H. Shirakawa, E.J. Louis, S.C. Gau, and A.G. MacDiarmid. 1977. Electrical conductivity in doped polyacetylene. Phys Rev Lett 39 (17):1098.
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129. Haupt, K., and K. Mosbach. 2000. Molecularly imprinted polymers and their use in biomimetic sensors. Chem Rev 100:2495. 130. Yu, J.C.C., and E.P.C. Lai. 2005. Interaction of ochratoxin A with molecularly imprinted polypyrrole film on surface plasmon resonance sensor. React Funct Polym 63:171. 131. Hutchins, R.S., and L.G. Bachas. 1995. Nitrate-selective electrode developed by electrochemically mediated imprinting=doping polypyrrole. Anal Chem 67:1654. 132. Yeh, W.-M., and K.-C. Ho. 2005. Amperometric morphine sensing using a molecularly imprinted polymer-modified electrode. Anal Chim Acta 542:76. 133. Campbell, T.E., A.J. Hodgson, and G.G. Wallace. 1999. Incorporation of erythrocytes into polypyrrole to form the basis of a biosensor to screen for rhesus (D) blood groups and rhesus (D) antibodies. Electroanalysis 11 (4):215. 134. Haddour, N., S. Cosnier, and C. Gondran. 2004. Electrogeneration of a biotinylated poly(pyrrole– ruthineum (II)) film for the construction of photoelectrochemical immunosensor. Chem Comm 2472. 135. Brahim, S., A.M. Wilson, D. Narinesingh, E. Iwuoha, and A. Guiseppi-Elie. 2003. Chemical and biological sensors based on electrochemical detection using conducting electroactive Polymers. Microchim Acta 143 (2–3):123. 136. Guiseppi-Elie, A., A.M. Wilson, and A.S. Sujdak. 1998. Electroconductive gels for controlled electrorelease of bioactive peptides. ACS Symp Ser 709, chap. 15, 185. 137. Misoska, V., J. Ding, J.M. Davey, W.E. Price, S.F. Ralph, and G.G. Wallace. 2001. Polypyrrole membranes containing chelating ligands: Synthesis, characterization and transport studies. Polymer 42:8571. 138. Atobe, M., S. Hitose, and T. Nonaka. 1999. Chemistry in centrifugal fields Part I. Electrooxidative polymerization of aniline. Electrochem Commun 1:278. 139. Eftekhari, A. 2004. Enhanced stability and conductivity of polypyrrole film prepared electrochemically in the presence of centrifugal forces. Synth Met 142:305. 140. Smela, E., G. Zuccarello, H. Kariis, and B. Liedberg. 1998. Thiol modified pyrrole monomers. Part 1. Synthesis, characterization, and polymerization of 1-(2-thioethyl)pyrrole and 3-(2-thioethyl)pyrrole. Langmuir 14 (11):2970. 141. Smela, E., H. Kariis, Z.P. Yang, K. Uvdal, G. Zuccarello, and B. Liedberg. 1998. Thiol modified pyrrole monomers. Part 2. As-deposited monolayers of 1-(2-thioethyl)pyrrole and 3-(2-thioethyl)pyrrole. Langmuir 14 (11):2976. 142. Smela, E., H. Kariis, Z.P. Yang, M. Mecklenburg, and B. Liedberg. 1998. Thiol modified pyrrole monomers. Part 3. Electrochemistry of 1-(2-thioethyl)pyrrole and 3-(2-thioethyl)pyrrole monolayers in propylene carbonate. Langmuir 14 (11):2984. 143. Smela, E. 1998. Thiol modified pyrrole monomers. Part 4. Electrochemical deposition of polypyrrole over 1-(2-thioethyl)pyrrole. Langmuir 14 (11):2996. 144. Idla, K., O. Inganas, and M. Strandberg. 2000. Good adhesion between chemically oxidized titanium and electrochemically deposited polypyrrole. Electrochim Acta 45 (13):2121. 145. Guiseppi-Elie, A., A.M. Wilson, J.M. Tour, T.W. Brockmann, P. Zhang, and D.L. Allara. 1995. Specific immobilization of electropolymerized polypyrrole thin films onto interdigitated microsensor electrode arrays. Langmuir 11:1768. 146. Abraham, S., S. Brahim, K. Ishihara, and A. Guiseppi-Elie. 2005. Molecularly engineered p(HEMA)-based hydrogels for implant biochip biocompatibility. Biomaterials 26:4767. 147. Bongiovanni, C., T. Ferri, A. Poscia, M. Varalli, R. Santucci, and A. Desideri. 2001. An electrochemical multienzymatic biosensor for determination of cholesterol. Bioelectrochemistry 54:17. 148. Singh, S., A. Chaubey, and B.D. Malhotra. 2004. Amperometric cholesterol biosensor based on immobilized cholesterol esterase and cholesterol oxidase on conducting polypyrrole films. Anal Chim Acta 502 (2):229. 149. Yamato, H., M. Ohawa, and W. Wernet. 1995. A polypyrrole=three enzyme electrode for creatinine detection. Anal Chem 67:2776.
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169. Soloducho, J. 1999. Convenient synthesis of polybispyrrole system. Synth Met 99 (3):181. 170. Jerome, C., V. Geskin, R. Lazzaroni, J.L. Bredas, A. Thibaut, C. Calberg, I. Bodart, M. Mertens, L. Martinot, D.J. Rodrigue Riga, and R. Jerome. 2001. Full-electrochemical preparation of conducting=insulating binary polymer films. Chem Mater 13:1656. 171. Cosnier, S. 1999. Biomolecule immobilization on electrode surfaces by entrapment or attachment to electrochemically polymerized films. A review. Biosens Bioelectron 14 (5):443. 172. Mailley, P. 2001. Inclusion of metal micro-particles into poly(pyrrolylalkylammonium) films containing an enzyme for bioelectrocatalysis: A preliminary study. Talanta 55 (5):1005. 173. Cosnier, S., F. Lambert, and M. Stoytcheva. 2000. A composite clay glucose biosensor based on an electrically connected HRP. Electroanalysis 12 (5):356. 174. Leca, B., R.M. Morelis, and P.R. Coulet. 1994. Influence of phospholipidic microenvironment on the performance of a polypyrrole enzyme electrode. Talanta 41 (6):925. 175. Yabuki, S., F. Mizutani, and M. Asai. 1991. Preparation and characterization of an electroconductive membrane containing glutamate dehydrogenase, NADP, and mediator. Biosens Bioelectron 6 (4):311. 176. Khan, G.F., E. Kobatake, H. Shinohara, Y. Ikariyama, and M. Aizawa. 1992. Molecular interface for an activity controlled enzyme electrode and its application for the determination of fructose. Anal Chem 64 (11):1254. 177. Aizawa, M., T. Haruyama, G.F. Khan, E. Kobatane, and Y. Ikariyama. 1994. Electronically modulated biological functions of molecular interfaced enzymes and living cells. Biosens Bioelectron 9 (9):601. 178. Koopal, C.G.J., and R.J.M. Nolte. 1994. Kinetic study of the performance of third-generation biosensors. Bioelectrochem Bioenerg 33 (1):45. 179. Kuwabata, S., and C.R. Martin. 1994. Mechanism of the amperometric response of a proposed glucose sensor based on a polypyrrole-tubule-impregnated membrane. Anal Chem 66 (17):2757. 180. Small, C.J., C.O. Too, and G.G. Wallace. 1997. Responsive conducting polymer-hydrogel composites. Polym Gels Networks 5:251. 181. Asberg, P., and O. Inganas. 2003. Hydrogels of a conducting conjugated polymer as 3-D enzyme electrode. Biosens Bioelectron 19:199. 182. Pich, A., Y. Lu, H.-J.P. Adler, T. Schmidt, and K.-F. Arndt. 2002. Dispersion polymerization of pyrrole in the presence of poly(vinyl methyl ether) microgels. Polymer 43:5723. 183. Douglass, P.M., S. Daunert, J.D. Patel, L.G. Bachas, K.-Q. He, and M.J. Madou. 2000. Biologically inspired, intelligent muscle material for sensing and responsive delivery of countermeasures. Soc Automot Eng 1. 184. Brahim, S., D. Narinesingh, and A. Guiseppi-Elie. 2002. Polypyrrole-hydrogel composites for the construction of clinically important biosensors. Biosens Bioelectron 17:53. 185. Rubio Retama, J., E. Lopez Cabarcos, D. Mecerreyes, and B. L’opez-Ruiz. 2004. Design of an amperometric biosensor using polypyrrole-microgel composites containing glucose oxidase. Biosens Bioelectron 20:1111. 186. Nishizawa, M., and I. Uchida. 1999. Microelectrode-based characterization systems for advanced materials in battery and sensor applications. Electrochim Acta 44:3629. 187. Matsue, T., M. Nishizawa, T. Sawaguchi, and T. Uchida. 1991. An enzyme switch sensitive to NADH. J Chem Soc Chem Commun 15:1029. 188. Bartlett, P.N., and P.R. Birkin. 1994. A microelectrochemical enzyme transistor responsive to glucose. Anal Chem 66:1552. 189. Bartlett, P.N., J.H. Wang, and E.N.K. Wallance. 1996. A microelectrochemical switch responsive to NADH. Chem Commun 3:359. 190. Bartlett, P.N, and P.R. Birkin. 1993. Enzyme switch responsive to glucose. Anal Chem 65:1118. 191. Barton, A.C., S.D. Collyer, F. Davis, D.D. Gornall, K.A. Law, E.C.D. Lawrence, D.W. Mills, S. Myler, J.A. Pritchard, M. Thompson, and S.P.J. Higson. 2004. Sonochemically fabricated microelectrode arrays for biosensors offering widespread applicability: Part I. Biosens Bioelectron 20:328.
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192. Lee, C.S., J. Joo, S. Han, and S.K. Koh. 2005. An approach to durable PVDF cantilevers with highly conducting PEDOT=PSS (DMSO) electrodes. Sens Actuators A 373. 193. Mattioli-Belmonte, M., G. Giavaresi, G. Biagini, L. Virgili, M. Giacomini, M. Fini, F. Giantomassi, D. Natali, P. Torricelli, and R. Giardino. 2003. Tailoring biomaterial compatibility: In vivo tissue response versus in vitro cell behavior. Int J Artif Organs 26:1077. 194. Smela, E. 2003. Conjugated polymer actuators for biomedical applications. Adv Mater 15 (6):481. 195. Schmidt, C.E., V.R. Shastri, J.P. Vacanti, and R. Langer. 1997. Stimulation of neurite outgrowth using an electrically conducting polymer. Proc Natl Acad Sci USA 94:8948. 196. Rivers, T.J., T.W. Hudson, and C.E. Schmidt. 2002. Synthesis of a novel, biodegradable, electrically conducting polymer for biomedical applications. Adv Funct Mater 12 (1):33. 197. Inoue, S., M. Yasuzawa, and S. Imai. 2002. Preparation of glucose sensor using organopolysiloxane=polypyrrole complex film. Proc 34th Chem Sens Symp 18 (A):136. 198. Persaud, K., and G.H. Dodd. 1982. Analysis of discrimination mechanisms in the mammalian olfactory system using a model nose. Nature 299:352. 199. Hodgins, D. 1995. The development of an electronic ‘nose’ for industrial and environmental applications. Sens Actuators B Chem 27 (1–3):255. 200. Nake, A., B. Dubreuil, C. Raynaud, and T. Talou. 2005. Outdoor in situ monitoring of volatile emissions from wastewater treatment plants with two portable technologies of electronic noses. Sens Actuators B Chem 106 (1):26. 201. Magan, N., A. Pavlou, and I. Chrysanthakis. 2001. Milk-sense: A volatile sensing system recognizes spoilage bacteria and yeasts in milk. Sens Actuators B Chem 72 (1):28. 202. Bourgeois, W., and R.M. Stuetz. 2002. Use of a chemical sensor array for detecting pollutants in domestic wastewater. Water Res 36 (18):4505. 203. Brezmes, J., Ma. L.L. Fructuoso, E. Llobet, X. Vilanova, I. Recasens, J. Orts, G. Saiz, and X. Correig. 2005. Evaluation of an electronic nose to assess fruit ripeness. IEEE Sens J 5 (1):97. 204. Guadarrama, A., J.A. Ferna´ndez, M. I´n˜iguez, J. Souto, and J.A. de Saja. 2001. Discrimination of wine aroma using an array of conducting polymer sensors in conjunction with solid-phase microextraction (SPME) technique. Sens Actuators B Chem 77 (1–2):401. 205. Guadarrama, A., M.L. Rodrı´guez-Me´ndez, C. Sanz, J.L. Rı´os, and J.A. de Saja. 2001. Electronic nose based on conducting polymers for the quality control of the olive oil aroma: Discrimination of quality, variety of olive and geographic origin. Anal Chim Acta 432 (2):283. 206. Fend, R., C. Bessant, A.J. Williams, and A.C. Woodman. 2004. Monitoring haemodialysis using electronic nose and chemometrics. Biosens Bioelectron 19:1581. 207. Lonergan, M.C., E.J. Severin, B.J. Doleman, S.A. Beaber, R.H. Grubbs, and N.S. Lewis. 1996. Arraybased vapor sensing using chemically sensitive, carbon black-polymer resistors. Chem Mater 8:2298. 208. Hanson, C.W. III, and E.R. Thaler. 2005. Electronic nose prediction of a clinical pneumonia score: Biosensors and microbes. Anesthesiology 102:63. 209. Ramanathan, K., M.A. Bangar, M. Yun, W. Chen, A. Mulchandani, and N.V. Myung. 2004. Individually addressable conducting polymer nanowires array. Nanoletters 4 (7):1237. 210. Gerard, M., A. Chaubey, and B.D. Malhotra. 2002. Application of conducting polymers to biosensors. Biosens Bioelectron 17:345. 211. Guiseppi-Elie, A., G.G. Wallace, and T. Matsue. 1998. Chemical and biological sensors based on electrically conducting polymers. In Handbook of Conducting Polymers, 2nd Edn, Eds. T. Skotheim, R. Elsenbaumer, and J.R. Reynolds, Chapter 34, pp. 963–991. New York: Marcel Dekker.
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13 Optical Biosensors Based on Conjugated Polymers 13.1 13.2 13.3
Introduction..................................................................... 13-1 Principles of Detection ................................................... 13-2 Biosensors Utilizing Quenching of Conjugated Polymers........................................................................... 13-4 Sensors for Genetic Material—DNA=RNA Sensors . Sensors for Ligand–Receptor Interactions . Sensors for Measuring Enzymatic Activity
13.4
Peter Nilsson and Olle Ingana¨s
13.1
Biosensors Utilizing Conformational Changes of Conjugated Polyelectrolytes ..................................... 13-11 Sensors for Genetic Material—DNA=RNA Sensors . Sensors for Ligand–Receptor Interactions . Sensors for Recording Conformational Changes in Proteins
Introduction
When considering new sensory technologies, one should look to nature for the most favorable solution. The evolution of living organisms has developed the ultimate chemical sensors. The astonishing sensory performance of biological systems does not originate from a single element. Their optimal performance is mostly derived from a completely interactive system wherein the selectivity is derived from receptors, and sensitivity is the result of analyte-triggered biochemical cascades. Clearly, optimal artificial sensory systems should also display all of these features. In this regard, conjugated polymers are a diverse sensor platform, and can be used in a wide range of biomolecular recognition schemes to obtain sensory responses. Biosensors based on conjugated polymers are sensitive to very minor perturbations, due to amplification by a collective system response and offer a key advantage compared to small molecules– based sensors [1–4]. The physical properties of conjugated polymers can be utilized for a wide range of biosensors [5] based on the mechanisms related to charge or energy transfer in its widest scope. The nature of conjugated polymers suggests biosensors based on charge transfer and transport, and electrochemical methods where the analyte controls the condition for charge transfer at electrodes have been studied [6–9]. We will leave these mechanisms out of discussion and here limit the coverage to the use of conjugated polymers as optical probes, and colorimetric biosensors and biosensors utilizing fluorescence will be discussed. The long history of such biosensors based on small molecules, either covalently attached to biological macromolecules or bound to them in solution, testifies to the need for sensors of this kind. The possibility to develop the staining of biomolecular structures, in order to follow morphology and dynamics with new probes, could lead to novel tools for biological research [10]. 13-1
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Receptor S
S
S S
S
n
S
S S
S S
n
FIGURE 13.1 Schematic drawing of the detection of a ligand–receptor interaction with a ligand (black triangles)functionalized polythiophene. The backbone of the polythiophene is undergoing a coil-to-rod (less conjugated to more conjugated) transition, upon ligand–receptor interactions.
The application of conjugated polymer for the colorimetric detection of biological targets (biochromism) was first reported by Charych in 1993 [11]. The technique is based on a ligand-functionalized conjugated polymer, which undergoes a colorimetric transition (coil-to-rod transition of the conjugated backbone), upon interaction with a receptor molecule (Figure 13.1). Analyte specificity in this first generation of conjugated polymer–based biosensors was due to the covalent integration of ligands on the side chains of the conjugated polymers. Ligand-functionalized versions of self-assembling polydiacetylenes have been used extensively for the colorimetric detection of ligand–receptor interactions and molecular interactions [11–15]. Polythiophene derivatives that display biotin [16–18] and different carbohydrates [19] have been synthesized and shown to undergo colorimetric transitions in response to binding of streptavidin and different types of bacteria and viruses, respectively. However, in all cases, the detection and recognition event is a function of the nature and characteristics of the side chains, and the side chain functionalization of the conjugated polymer requires advanced synthesis and extensive purification of numerous monomeric and polymeric derivatives. This first generation of sensors was also mainly using optical absorption as the source for detection, and the sensitivity of these sensors were much lower compared with other sensing systems for biological processes. However, there might be a revisit for this class of sensors if these problems are solved. A fluorescent poly(p-phenylene ethynylene) derivative with carbohydrate-functionalized side chains for the detection of different bacteria [20] and novel synthesis schemes for ligand functionalization of polythiophenes [21] were recently reported. To avoid covalent attachment of the receptor to the polymer side chain and to increase the sensitivity of the biosensors, fluorescent-conjugated polyelectrolytes have been utilized. Most of the presently demonstrated systems [4,22–25] use anionic-, cationic-, or zwitterionic-conjugated polyelectrolytes, with interactions dominated by electrostatic forces. Conjugated polythiophenes offer a new route for polymer–biopolymer interactions where hydrogen bonding, having a major influence on chain interactions, has also been reported [26–29]. In the following sections, biosensors based on fluorescentconjugated polyelectrolytes will be discussed. These sensors will be dealt with according to the principle used for the detection of the biospecific interaction.
13.2
Principles of Detection
The most sensitive biosensors based on conjugated polyelectrolytes reported in the literature are utilizing changes in the absorption or emission properties from the conjugated polyelectrolytes. As the conjugated polyelectrolyte’s absorption and emission characteristics are largely determined by the local electronic structure, the sensitivity of the band gap, and by implication the optical properties, to the backbone conformation of the conjugated polyelectrolyte provides a useful means to create this type of sensors. Fluorescence is a widely used and rapidly expanding method in chemical sensing and aside from inherent sensitivity, this method offers diverse transduction schemes based upon changes in intensity,
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Optical Biosensors Based on Conjugated Polymers
energy transfer, wavelength (excitation and emission), polarization, and lifetime. Optical sensors, based on conjugated polyelectrolytes, can mainly be divided into two different types, depending on which detection scheme is used. Schematic drawings of the two detection schemes for the detection of DNA hybridization are shown in Figure 13.2. In the first approach, quenching of the fluorescence from the conjugated polyelectrolyte chain is used [1–5,30–36]. The quenching may be due to fluorescence resonance energy transfer (FRET) to another emitter or due to (dark) excitation quenching. If the quencher is combined with a biomolecule in a covalently bonded object, and this object is coordinated to the conjugated polyelectrolyte chain by weak noncovalent interactions (electrostatic or hydrophobic interactions), it is possible to detect the presence of a certain biomolecule in a sample by the quenching of the emitted light from the conjugated polyelectrolyte. Removal of the quencher from the chain turns on luminescence; as the chain is long, it is possible to probe a large volume and, therefore, low concentrations of the target to be detected. The mechanism behind this very sensitive detection has been named superquenching and is still debated. A wide range of conjugated polyelectrolytes have been used as detecting elements for biological molecules in an aqueous environment, and many of these systems utilize the impact of biomolecules on the conditions for FRET or excitation transfer [4,24,37–52]. The second type of biosensors are based on the detection of biological processes through their impact on the conformation and the geometry of the polyelectrolyte chains [6,22,23,25,28,29,53–62]. The conformational flexibility of conjugated polyelectrolytes allows direct correlation between the geometry of chains and the resulting electronic structure and processes. To use this phenomenon as a sensor for the recording of conformational changes of biomolecules, the conjugated polymer chain geometry is required to be governed by the conformational changes of the biomolecules. If conformational changes of biomolecules can lead to different conformations of the polymer backbone, an alteration of the
+
Cationic polyfluorene
+
Chromophore/quencher-labeled single-stranded PNA probe
Helical zwitterionic polythiophene
Single-stranded DNA probe (ssDNA)
Hybridization Polyelectrolyte/sDNA complex Rod-shaped and aggregated Polyelectrolyte chains Energy transfer
Polyelectrolyte/dsDNA complex
Hybridization
Polyelectrolyte/dsDNA complex Nonplanar and separated polyelectrolyte chains
FIGURE 13.2 Schematic drawing of the detection of DNA hybridization by using conjugated polyelectrolytes. The technique using FRET or quenching from the CP chain (left) and the technique using conformational change upon interaction with biomolecules (right) are given.
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absorption and emission properties from the polymer will be observed. Hence, conjugated polyelectrolytes can be used as conformation-sensitive optical probes, appropriate for making novel biosensors. The technique has been used to detect conformational changes in synthetic peptides [53,59], calciuminduced conformational changes in calmodulin [60], and the formation of protein amyloid fibrils [28,61]. The detection of these biological processes is utilized by using the conformational changes of anionic and zwitterionic polythiophene derivatives, which are noncovalently attached to the biomolecule of interest. Conjugated polyelectrolytes can be used for creating biosensors for the global geometrical changes occurring as biomacromolecules recognize their kind, and as biomacromolecules change conformation. This is something quite different from what has been accomplished with many fluorescent detector dyes over the last 100 years. Rather than adding pointlike fluorophores by covalent chemistry to well-defined sites on a biological macromolecule, noncovalent complexes between a synthetic luminescent conjugated polyelectrolyte and a biological polyelectrolyte are formed. These complexes cause changes of geometry of the conjugated polyelectrolyte, and geometrical changes of this complex can then be followed in situ, by optical absorption and emission.
13.3
Biosensors Utilizing Quenching of Conjugated Polymers
Excitation quenching, also labeled superquenching, has been described in several reports and is based on the finding that photoluminescence of conjugated polyelectrolytes can be quenched by means of energy and electron transfer to small molecule quenchers, chromophores, or metals [1–3,5,23,31–36]. It has been claimed that one quencher molecule can quench the photoluminescence of up to several hundred polymer-repeating units [1–5,30–36]. The use of fluorescence quenching of conjugated polyelectrolytes has been utilized for a wide range of bioassays, including DNA hybridization, ligand–receptor interactions, and the measurement of enzymatic activity.
13.3.1 Sensors for Genetic Material—DNA=RNA Sensors The detection of DNA hybridization is of great interest for genetic analysis to diagnose bacterial and viral infectious diseases, associated with mutations in the DNA sequence. An effective DNA-sensing system requires selectivity to sense single-nucleotide polymorphisms (SNPs) and sensitivity to detect a small number of copies of DNA. It is also of great interest to implement the system in the form of a microarray, as multiple DNA sequences often has to be identified. Hence, an effective DNA-sensing system requires anchoring and patterning of the detecting element on a surface. DNA sensors based on quenching of conjugated polyelectrolyte (donor) with FRET to a smaller chromophore dye (acceptor) have been reported [24,40–43,45,46]. A schematic drawing of these sensors is shown in Figure 13.2 and Figure 13.3. The technique is based on noncovalent assembly of a cationic conjugated polyelectrolyte to chromophore-labeled DNA or peptide nucleic acid (PNA) molecules. When the hybridization takes place, a decrease in the average distance between the conjugated polyelectrolyte chain and the chromophore occur due to electrostatic interactions between the two molecules. By forming a complex between the cationic-conjugated polyelectrolyte and the chromophore-labeled molecule, FRET is allowed [24]. Interestingly, by using the conjugated polyelectrolyte as a light-harvesting donor, the emission from the acceptor was increased (>25 times) compared to direct excitation of the acceptor. This amplification allows the detection of DNA concentration as low as 10 pM (1011 M) with an ordinary fluorometer [24]. This technique has also been developed further and recently SNP detection [45] and incorporation of the method into DNA chips and microarrays [46] were reported. However, the SNP detection requires the addition of an enzyme (nuclease) and the chip technology requires the covalent attachment of the DNA=PNA molecule to the surface. This technique might be utilized to make sensitive and selective DNA chips based on conjugated polyelectrolytes. A similar technique, using fluorescent polymer quenching for DNA sensing, has also been reported [37,39]. This technique is more complex compared with the one described above and requires the use of other molecules than a conjugated polyelectrolyte and a quencher-labeled DNA or PNA molecule. The
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+ + +
+
+ + + Cationic conjugated polyelectrolyte (CCP)
Chromophore-labeled single-stranded PNA probe
− −
− Addition of DNA
−
−
−
FRET
+ +
+ + +
−
+
−
+ −
+ −
+
− Complementary
− − −
CCP/dsDNA complex
+ + +
Non complementary
− + −
− −
+ − + −
+
+ −
+
+ −
+
+ −
CCP/ssDNA complex
FIGURE 13.3 Schematic presentation of a detection scheme for DNA hybridization, utilizing quenching of a cationic conjugated polyelectrolyte by a fluorophore-labeled PNA probe.
basis of this assay is polystyrene microspheres that are coated with neutravidin (a biotin-binding protein) and a biotinylated anionic fluorescent conjugated polyelectrolyte. A biotinylated PNA or DNA strand serves as a capture ligand for DNA nucleotides and when mixed with the microspheres, a strong complex is created through the biotin–avidin interaction [37,39]. The assay can then be used in three different ways (Figure 13.4): 1. Microspheres with a capture ligand are presented to a quencher-labeled DNA strand. Quenching occurs if this strand is complementary and binds to the capture ligand. This is a direct quenching assay. 2. The microspheres are first presented to an unlabeled DNA strand that is complementary to the capture ligand. A fixed concentration of the quencher-labeled DNA strand is then added and quenching occurs. The strength of this quenching is dependent on the amount of unlabeled DNA strand that has bound. This is a competitive assay.
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Addition of quencherlabeled DNA MS with anionic CP and capture-ligand DNA
Addition of unlabeled DNA
Strong quenching
Addition of quencherlabeled DNA
Unlabeled DNA depending quenching
MS with anionic CP and capture-ligand DNA
Addition MS with anionic CP
Mixture of capture-ligand DNA, unlabeled DNA, and quencherlabeled DNA in solution
Unlabeled DNA depending quenching
FIGURE 13.4 A quenching assay where microspheres (MS) coated with an anionic conjugated polyelectrolyte (CP) and capture-ligand DNA are presented to quencher-labeled DNA. Quenching is effected upon interaction between the capture-ligand DNA and the chromophore-labeled DNA (top). A competitive quenching assay where MS coated with an anionic CP and capture-ligand DNA are presented to unlabeled DNA. Quencher-labeled DNA strands are then added and quenching occurs. The strength of this quenching is dependent on the amount of unlabeled DNA strand that has bound (middle). A solution competitive assay where capture-ligand DNA, unlabeled DNA, and quencher-labeled DNA are mixed in solution. The solution is then mixed with MS coated with an anionic CP and quenching occurs. The strength of this quenching is dependent on the amount of unlabeled DNA strand that has bound (bottom).
3. A fixed amount of the quencher-labeled DNA strand and variable amounts of an unlabeled DNA strand are mixed with the capture ligand in solution. The reaction mixture is then mixed with the conjugated polyelectrolyte-coated microspheres and the quenching that is observed depends on the amount of unlabeled DNA strand that is present. This is also a competitive assay.
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These assays have been used to detect the presence of subpicomolar concentrations of DNA in solution and for the detection of SNPs [37,39].
13.3.2 Sensors for Ligand–Receptor Interactions Similar techniques, as described in the previous section, have been used to detect ligand–receptor interactions [4,38,50,63]. Chen et al. [4] demonstrated that the luminescence of an anionic conjugated polyelectrolyte could be quenched by extremely low concentrations of a cationic electron acceptor molecule, methylviologen (MV2þ). The huge enhancement of the fluorescence quenching arises from a combination of two effects. First, the local concentration of the quencher is enhanced, due to the formation of a complex between the negatively charged polyelectrolyte and the cationic electron acceptor. Second, it appears that a single MV2þ molecule can quench the fluorescence from an entire polyelectrolyte chain. The mechanism of this remarkable effect is not yet understood, but it has been observed in other systems [2–5,30–36]. The technique has so for been demonstrated for biotin–avidin interactions [4,50] and a quencher– antibody interaction [63] (Figure 13.5). The quencher molecules are covalently attached to the ligand (biotin) and by addition of the receptor (avidin), the quencher is removed or attached to the conjugated polyelectrolyte. As the interactions between avidin and the conjugated polyelectrolyte depend on the pH of the systems, both the addition and the removal of the quenching molecule are possible [50]. Removal of the quencher results in an enhanced fluorescence from the conjugated polyelectrolyte and attachment of the quencher can be seen as a decrease of the fluorescence from the conjugated polyelectrolyte. This technique has the ability to create sensitive biosensors for a diverse range of ligand–receptor interactions [38]. However, covalent attachment of a quencher to the ligand is necessary and there are also problems with unspecific binding of the conjugated polyelectrolyte to the receptor molecule [50].
13.3.3 Sensors for Measuring Enzymatic Activity Fluorescent polyelectrolyte quenching assays for a host of kinases, phosphatases, and proteases have been reported [48,49,51,52]. The first reported assay [48] is based on electrostatic interactions between an anionic conjugated polyelectrolyte and a peptide substrate that are labeled with a fluorescence quencher. Protease activity is measured in real time by using fluorescence spectroscopy and two approaches are presented (Figure 13.6). In the first approach, the peptide substrate is cleaved by the enzyme, whereby the quencher is released from the conjugated polyelectrolyte, resulting in a turn-on of the fluorescence from the conjugated polyelectrolyte. This turn-on system was used to sense enzyme activity when the concentrations of the enzyme and substrate are in the nanomolar regime [48]. In the second approach, peptide cleavage results in a turn-off of the fluorescence from the conjugated polyelectrolyte due to better interaction between the quencher and the polyelectrolyte upon hydrolysis of the peptide. Similar assays for measuring protease activity were also reported by Whitten and coworkers [49]. These assays utilize polystyrene microspheres that are coated with streptavidin (a biotin-binding protein) and a biotinylated anionic fluorescent conjugated polyelectrolyte or a cationic polyelectrolyte. A biotinylated quencher-labeled peptide serves as a substrate for the enzyme and when mixed with the microspheres, a strong complex similar to the one in the DNA assays described in Section 13.3.1, through the biotin–avidin interaction is created [49]. A schematic drawing of the assay is shown in Figure 13.7. Microspheres coated with streptavidin and biotinylated anionic-conjugated electrolytes are presented to a quencher-labeled biotinylated peptide substrate. Quenching of the fluorescence from the conjugated polyelectrolyte occurs as the peptide binds to the microsphere. The quencher is removed when the peptide is cleaved by the protease, detected as enhancement of the fluorescence from the conjugated polyelectrolyte. The assay can also be used with a cationic-conjugated polyelectrolyte and solution sensors based on a polyelectrolyte–avidin ensemble were also reported [49]. These
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Avidin Q+
Q+
+
Unquenched anionic conjugated polyelectrolyte
Quenched anionic conjugated polyelectrolyte
Avidin +
Q
+
Q+
Unquenched anionic conjugated polyelectrolyte
Quenched anionic conjugated polyelectrolyte
Antibody
Q+
Quenched anionic conjugated polyelectrolyte
+ Q+
Unquenched anionic conjugated polyelectrolyte
FIGURE 13.5 Schematic representation of conjugated polyelectrolyte sensors for ligand–receptor interactions. (Top) Quencher (Q)–Biotin (black circle) is removed by avidin (white square) and results in unquenched polyelectrolyte (dotted line). (From Chen, L.H., McBranch, D.W., Wang, H.L., Helgeson, R., Wudl, F., and Whitten, D.G., Proc. Natl. Acad. Sci. USA, 96, 12287, 1999.) (Middle) Quencher (Q)–Biotin (black circle) is added to conjugated polyelectrolyte by avidin (white square) and results in quenched polyelectrolyte (dotted line). (From Dwight, S.J., Gaylord, B.S., Hong, J.W., and Bazan, G.C., J. Am. Chem. Soc., 126, 16850, 2004.) (Bottom) The quencher (Q) is removed by quencher-specific antibody and results in unquenched polyelectrolyte (dotted line). (From Wang, D.L., Gong, X., Heeger, P.S., Rininsland, F., Bazan, G.C., and Heeger, A.J., Proc. Natl. Acad. Sci. USA, 99, 49, 2002.)
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Protease
+
+
+
Quenched anionic conjugated polyelectrolyte
Unquenched anionic conjugated polyelectrolyte
+
Protease
+
+
+
+
Unquenched anionic conjugated polyelectrolyte
Quenched anionic conjugated polyelectrolyte
FIGURE 13.6 Mechanism of the ‘‘turn-on’’ (top) and ‘‘turn-off ’’ (bottom) proteases sensors based on quenching of a conjugated polyelectrolyte (dotted line) by a quencher (black circle)-labeled peptide (black line).
fluorescent polymer–quenching assays for protease activity offer advantages over antibody-based assays, in terms of simplicity and sensitivity. It can also be used for a lot of different proteases and has the required specificity for screening chemical libraries for novel inhibitors of protease activity in a high-throughput screening (HTS) environment. The assay is also robust and platform-independent. It can be used in single test tubes, 96-well, and 384-well plate format [49].
B
Protease
Peptide substrate labeled with biotin and quencher Microsphere coated with streptavidin and a conjugated polyelectrolyte
B
Quenched conjugated polyelectrolyte
B
+
Unquenched conjugated polyelectrolyte
FIGURE 13.7 General scheme for a microsphere protease assay. The fluorescence of the conjugated polyelectrolyte is enhanced as the quencher is removed by cleavage of the peptide.
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Kinase PO3− Phosphatase Fluorescence “ON”
Fluorescence “OFF ”
PO3−
Microsphere coated with Ga3+ and an anionicconjugated polyelectrolyte
Microsphere coated with Ga3+ and an anionicconjugated polyelectrolyte
FIGURE 13.8 General scheme for a microsphere kinase=phosphates assay. The fluorescence of the conjugated polyelectrolyte is quenched as the quencher-labeled peptide (black line with a filled black circle) is brought in close vicinity to the microsphere, due to interaction between Ga3þ and PO3.
The same research group has also developed an assay for the detection of kinase and phosphate activities [51]. Kinases and phosphatases are enzymes responsible for phosphorylation and dephosphorylation of proteins, which regulates cellular metabolism, growth, differentiation, and proliferation. The assay is based on metal ion–mediated conjugated polyelectrolyte quenching (Figure 13.8). Anionicconjugated polyelectrolytes were used to coat quaternary amine latex spheres and these spheres were additionally coated with Ga3þ ions. The Ga ions serves as capturing agents for the phosphate group on peptide substrates. The microspheres are presented to different quencher-labeled peptides and the peptides bind to the microsphere depending on the phosphorylation of the peptide. If binding occurs, the fluorescence from the conjugated polyelectrolyte is decreased. The assay has been used for a wide range of different kinases or phosphatases and the modulation of the fluorescence signal from the conjugated polyelectrolyte is proportional to the enzyme activity [51]. It can be used for HTS of kinase or phosphate inhibitors and is a valuable tool for drug discovery. Recently another approach for the detection of protease activity by using fluorescent quenching of a conjugated polyelectrolyte was reported [52] (Figure 13.9). This assay utilizes a covalent attachment of the quencher-labeled peptide substrate to the anionic-conjugated polyelectrolyte and combines the
Protease +
Quenched conjugated polyelectrolyte due to functionalization with a quencherlabeled peptide
Unquenched conjugated polyelectrolyte
FIGURE 13.9 Schematic drawing for protease assay containing a conjugated polyelectrolyte (dotted line) functionalized with a peptide substrate (black line) containing a quencher (filled black circle). The fluorescence of the conjugated polyelectrolyte is enhanced as the quencher is removed by cleavage of the peptide.
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approach seen for the first generation of conjugated polymer–based biosensors with the quenching phenomenon. By covalent attachment of the peptide to the side chain of the polyelectrolyte, an amplified turn-on fluorogenic probe for proteases is generated. The system shows a fluorescence enhancement of one order of magnitude, as the quencher is removed by treating the substrate with the enzyme.
13.4
Biosensors Utilizing Conformational Changes of Conjugated Polyelectrolytes
Conjugated polymers display a remarkable array of color transitions due to conformational changes of the polymer backbone. The color changes can be ascribed to a change in the effective conjugation length of the delocalized p-conjugated polymer backbone, as the conformational changes occur [64]. The type of biosensors described in this section are based on a concept that utilizes geometrical changes of water-soluble, ionic-conjugated polyelectrolytes, which can specifically transduce biological processes such as DNA hybridization or ligand–receptor interactions, into a clear colorimetric or fluorometric signal. Compared to the first colorimetric sensors [11–21] based on conjugated polymers and the earlier described quenching technique, this approach does not require any chemical modification or labeling of the receptors or the ligands. This new, simple, rapid, sensitive, and selective methodology is based on noncovalent interaction between the conjugated polyelectrolyte and the biomolecule of interest. The technique has so far been used for the detection of DNA hybridization [23,25,29,56,57,62], ligand–receptor interactions [6,25,54–56], and conformational changes in synthetic peptides and proteins [28,53,59–61].
13.4.1 Sensors for Genetic Material—DNA=RNA Sensors An assay, using polycationic-conjugated polyelectrolytes, not requiring chemical labeling of the target nucleic acid (receptor) or covalent attachment of the DNA probe (ligand), was recently reported [23] (Figure 13.10). By simply altering the geometry of the conjugated backbone, a high sensitivity (zeptomole) [57] and selectivity (detection of SNPs) [23,57] were reported for this method using fluorescence as the source for detection. The method is based on noncovalent interaction between a cationicconjugated polyelectrolyte and a single-stranded DNA (ssDNA) probe, forming a strong electrostatic complex. The formation of this complex will induce a planarization of the backbone of the conjugated polyelectrolyte and aggregation of polyelectrolyte chains, seen as a red shift of the absorption and emission maximum, and a decrease of the intensity of the emitted light. When finding the proper target DNA=RNA, complementary to the ssDNA probe, the backbone of the conjugated polyelectrolyte becomes more twisted and the polyelectrolytes chains become more separated. The DNA hybridization event is detected as a blue shift of the absorption maximum or as an enhanced fluorescence intensity [23,57]. However, the above mentioned assay method [23,57] requires DNA denaturation conditions for SNP detection. The multiplicity of DNA sequences to identify also requires that the method can be implemented in the form of a microarray, and thus requires anchoring and patterning of the detecting polymer on a surface. A similar method for DNA hybridization was recently reported [29]. This method is based on the same principle as described above, but utilizes conformational changes of a helical zwitterionic conjugated polyelectrolyte [26] instead of a cationic-conjugated polyelectrolyte. This zwitterionic-conjugated polyelectrolyte can form a strong electrostatic complex with ssDNA and in addition create versatile hydrogen-bonding patterns with both ssDNA and double-stranded DNA (dsDNA). The method is highly sequence-specific and an SNP can be detected within 5 min without using any denaturation steps [29]. The interaction with DNA and the optical phenomena persists when the polyelectrolyte is deposited and patterned on a surface, which offers a novel way to create DNA chips without using covalent attachment of the receptor or labeling of the analyte.
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+ Single - stranded DNA probe (ssDNA)
Conjugated polyelectrolyte (CP)
CP/ssDNA complex Rod-shaped and aggregated polyelectrolyte chains Decreased fluorescence
Complementary
Addition of DNA/RNA
+
Non complementary
CP/ssDNA complex Non planar and separated polyelectrolyte chains Enhanced fluorescence
CP/ssDNA complex Rod-shaped and aggregated polyelectrolyte chains Decreased fluorescence
FIGURE 13.10 General detection scheme for a DNA sensor, utilizing conformational changes of conjugated polyelectrolytes.
Recently a method using a fluorescent cationic polymer and PNA probes for the detection of DNA targets on solid support was reported [62]. The technique is utilizing a similar detection scheme as shown in Figure 13.10, but the ssDNA probe is replaced by a PNA probe covalently attached to a solid support. The cationic polymer does not bind to the neutral PNA probes, but strongly interacts with the negatively charged backbone of a complementary oligonucleotide bound to the PNA probes, allowing transduction of the hybridization event into a fluorescence signal. This technique offers a simple sensitive electrostatic approach on the one hand, which enables detection and specific detection of unlabeled target DNA analyte using a microarray scanner. However, the technique requires covalent attachment of the PNA probes to the solid support and DNA-capturing probes could not be used. On the other hand, the detection scheme described in this study [62] reported an increased sensitivity, 2.5 1013 mole of DNA in a volume of 20 mL, compared to the study [29] described above. The method can also be used for the detection of SNP at room temperature, so no dehybridizaton steps are necessary [62].
13.4.2 Sensors for Ligand–Receptor Interactions The technique of using conformational changes in conjugated polyelectrolytes has also been used for the detection of ligand–receptor interactions [6,54–56,60]. The first method was reported by Faid and Leclerc in 1998 [6], and this method was based on acid–base complexation between an anionicconjugated polyelectrolyte and amine-functionalized biotin [6]. When exposing this complex to the biotin-binding protein avidin, a red shift of the absorption maximum from the conjugated polyelectrolyte due to planarization of the polyelectrolyte backbone was observed [6,54]. The biotin–avidin interaction also changed the electrochemical properties of the conjugated polyelectrolyte, thus the electronic detection was also possible with this method. Interestingly, femtomoles (1015 moles) of the protein in an aqueous environment could be detected [55].
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α-Thrombin
Addition of CCP Quadruplex structure of the DNA aptamer
Non planar and separated polyelectrolyte chains Enhanced fluorescence
Single stranded DNA-aptamer
Addition of +
+
CCP Non specific protein Rod-shaped and aggregated polyelectrolyte chains Decreased fluorescence
FIGURE 13.11 Schematic description of the specific detection of a-thrombin by use of a ssDNA aptamer (black line) and a cationic-conjugated polyelectrolyte (black dotted line). (From Ho, H.A., and Leclerc, M., J. Am. Chem. Soc., 126, 1384, 2004.)
The same system used for the detection of DNA=RNA [23] has also been used for the detection of ligand–protein interactions [25,56]. In this case, the ssDNA function as an aptamer (ligand) that changes its conformation upon binding to different proteins. The presently reported system [56] demonstrates the use of a cationic-conjugated polyelectrolyte for the detection of such conformational changes of a ssDNA aptamer upon binding to human a-thrombin (Figure 13.11). If the conjugated polyelectrolyte is binding to the free-ssDNA aptamer, a red shift of the absorption maximum and a decrease of the intensity of the emitted light from the polyelectrolyte are observed due to planarization and aggregation of the polyelectrolyte chains. The conformational change of the ssDNA aptamer upon binding to a specific protein is recorded as a blue shift of the absorption maximum or an enhancement of the fluorescence from the conjugated polyelectrolyte due to twisting and separation of the polyelectrolyte chains [56]. By this highly sensitive and selective method, subpicomolar of a specific protein can be detected [56], and the procedure can be used to detect a wide range of proteins. The same concept has also been used for specific detection of potassium ions [56] and for enantiomeric resolution of nucleic acids [25], and might be used for identification of a wide range of molecules as well as for HTS for drug discovery. Conformational changes of a helical zwitterionic-conjugated polyelectrolyte, poly(3-((S)-5-amino-5carboxyl-3-oxapentyl)-2,5-thiophene) hydrochloride (POWT) [26], have also been used to detect ligand–receptor interactions [60]. This methodology is also based on noncovalent assembly between the conjugated polyelectrolyte and a protein, calmodulin (CaM). Calmodulin is a small protein (148 amino acids) that functions as the primary intracellular calcium sensor in eukaryotic cells and plays an essential role in calcium-mediated signal transduction [65]. The POWT–CaM complex can be used to evaluate such interactions by exposing the complex to calcineurin, a 77 kDa CaM-binding protein. The technology utilizes geometrical changes of the polyelectrolyte chains, which are associated with intraand interchain events. These events have also been observed in thin films of POWT [66] and will reduce
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the fluorescence quantum yield from the single-chain level of 26% (observed in dilute methanol solutions) to 4% in solid solutions, due to nonradiative de-excitation. Quantum chemical modeling gave evidence for the creation of a new interchain transition as oligothiophene models approach each other. This new channel for de-excitation is created in the contact between polyelectrolyte chains and the emission maximum for the intrachain and the interchain event in thin polyelectrolyte films was 565 and 670 nm, respectively [66]. The ratio of the intensity of the emitted light at 540 and 670 nm, 540=670 nm, can be used as a measurement of the geometrical alteration of the polyelectrolyte chains and has previously been used for the detection of biospecific interaction [53,59]. The mechanism of the system and the titration curve for POWT:CaM with different amount of calcineurin is shown in Figure 13.12. The ratio of the emitted light at 540=670 nm is altered with an increasing amount of calcineurin and the dissociation constant (KD) for CaM and calcineurin can be estimated to approximately 36 nM. The ratio of the intensity of the emitted light at 540=670 nm for the POWT–CaM complex is not altered when Ca2þ is not present in the solution. This result is expected, as calcium activation of CaM is needed for the interaction between CaM and calcineurin. Hence the alteration of the ratio of the intensity of emitted light at 540=670 nm is most likely due to the interaction between calcium-activated CaM and calcineurin. This interaction is probably leading to a separation of the polyelectrolyte chains seen as an enhanced ratio of the intensity of the emitted light at 540=670 nm. As a control, the calcium-activated POWT–CaM complex was also exposed to human serum albumin (HSA) and interestingly no significant change in the ratio of the emitted light at 540=670 nm could be seen (Figure 13.12). This result suggests that the nonspecific binding between the POWT–CaM complex and HSA is minimal and argues that it is the calcium-activated CaM that is interacting in a selective way with calcineurin.
13.4.3 Sensors for Recording Conformational Changes in Proteins The geometrical changes of conjugated polyelectrolytes have also been used for the recording of conformational changes in synthetic peptides [53,59] and proteins [28,60,61]. The same study [60], as described above, also showed how conformational alterations of a conjugated polyelectrolyte could be utilized to detect conformational changes in CaM. The overall structure of CaM (Figure 13.13) consists of two globular calcium-binding domains, each containing two calcium-binding regions with the characteristic EF hands [67], connected by a linker. Upon binding of calcium, the relative orientation of the two a-helices that define the EF-hand changes substantially, resulting in a transition from a closed to an open conformation of the protein motif. Structural studies have also shown that calcium activation of CaM is accompanied by a global conformational change, whereby the compact calcium-free form of CaM is converted to a more extended dumbbell-shaped molecule upon binding of calcium [68,69]. The extended form of the protein consists of two lobes separated by a central a-helix, and this central helix is flexible and allows considerable movements of the two lobes with respect to one another. By formation of a complex between POWT and CaM, the emission maximum of POWT is red-shifted and the intensity of the emitted light is decreased (Figure 13.13), indicating that the POWT backbone becomes more planar and that aggregation of the POWT chains occurs.By an addition of 10 mM Ca2þ to this complex, the emission maximum (594 nm) is blue-shifted and the shoulder around 540 nm is increased (Figure 13.13), suggesting that the polyelectrolyte backbone becomes more nonplanar and that a separation of the polymer chains occur. The ratio of the intensity of the emitted light at 540=670 nm is increased, showing that the conformational changes of the CaM molecule upon exposure to Ca2þ are governing the geometry of the polyelectrolyte chains. The increased intrachain event at 540 nm, associated with separation of the polyelectrolyte chains, are probably a result of the conformational change whereby the compact calcium-free form of CaM is converted to a more extended dumbbellshaped molecule upon binding of calcium [68,69]. A schematic presentation of the different conformational alterations of the CaM molecule upon exposure to calcium and the suggested POWT chains geometries seen for the different POWT=CaM solutions is shown in Figure 13.13 [60].
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Calcineurin Ca2+
Separation of POWT chains. Enhanced ratio 540/670 nm
+
POWT
CaM
POWT–CaM complex
ca2+
No separation of POWT chains. Unchanged ratio 540/670 nm
HSA
Difference in ratio 540/670 nm
0.50 0.45 0.40 0.35 0.30 0.25 0.20 0.15 0.10 0.05 0.00 0
50
100
150
Concentration (nM)
FIGURE 13.12 (Top) Proposed mechanism for the detection of calmodulin (CaM)–calcineurin interactions with POWT. (Bottom) The difference in ratio of the intensity of the emitted light (540=670 nm) for a POWT–CaM complex upon exposure to different amounts of calcineurin in 20 mM Tris–HCl, pH 7.5, 10 mM Ca2þ (squares); different amounts of calcineurin in 20 mM Tris–HCl, pH 7.4 (diamonds); and different amount of HSA in 20 mM Tris–HCl, pH 7.4, 10 mM Ca2þ (triangles).
POWT has also been used for the detection of conformational changes in synthetic peptides [59]. These peptides, one cationic and one anionic, were designed to adopt random-coil formations by themselves and when mixing the two peptides, heterodimers with a four-helix bundle conformation were formed (Figure 13.14). When mixing POWT with different forms of the peptides, alterations of the emission spectrum of POWT were observed (Figure 13.14) [59]. The addition of a positively charged peptide with a random-coil conformation, JR2K, will force the polyelectrolyte to adopt a nonplanar conformation with separated polyelectrolyte chains, observed as a blue shift and an increased intensity of the emitted light. Upon exposure to a negatively charged
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4.00E+05
Fluorescence (cps)
3.50E+05 3.00E+05 2.50E+05 2.00E+05 1.50E+05 1.00E+05 5.00E+04 0.00E+00 500
550
600
650
700
Wavelength (nm)
Ca2+
Ca2+
FIGURE 13.13 (See color insert following page 8-22.) (Top) Structure of calmodulin (CaM) without Ca2þ (left) and with Ca2þ (right). (Middle) Emission spectra of POWT (black line), POWT-CaM (red line), and POWT–CaM–Ca2þ (blue line) in 20 mM Tris–HCl, pH 7.5. (Bottom) Schematic drawing of the different conformational changes of the CaM molecule (blue helices) upon exposures to Ca2þ and the suggested geometries of the POWT chains (pink helices).
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JR2E Random coil
JR2K Random coil
Four-helix bundle (dimer)
4.00E+05 3.50E+05
+
Fluorescence (cps)
3.00E+05 2.50E+05
+
2.00E+05 1.50E+05 1.00E+05
+
5.00E+04 0.00E+00 500
550
600
650
700
Wavelength (nm)
FIGURE 13.14 (See color insert following page 8-22.) (Top) Schematic drawing of the conformation of the synthetic peptides JR2E (negatively charged) and JR2K (positively charged). (Bottom) Fluorescence spectra of POWT–JR2E (red line), POWT–JR2K (blue line), and POWT–JR2E–JR2K (green line).
peptide with a random-coil conformation, JR2E, the backbone adopts a planar conformation and aggregation of the polyelectrolyte chains occurs, seen as a red shift and a decreased intensity of the emitted light. By adding JR2K to the POWT–JR2E, the intensity of the emitted light is increased and blue-shifted, associated with separation of the polyelectrolyte chains. This geometrical alteration of the polyelectrolyte chains is due to the formation of a four-helix bundle of the peptides. Hence, different emission spectrum of POWT is seen, depending on the charge and conformational change of the peptides. Simple, sensitive, and versatile tools that detect the conformational changes in proteins are of great importance, as many diseases are associated with conformational changes in proteins. The importance of conformational changes of proteins leading to pathogenic states, such as Alzheimer’s disease, the systemic amyloidosis, and transmissible spongiform encephalopathy (TSE), has been well documented [70–73]. Especially under conditions that destabilize the native state, proteins can aggregate into characteristic fibrillar assemblies, known as amyloid fibrils [72] (Figure 13.15). Many proteins are readily converted to an inactive b-sheet-rich amyloid fibrillar form by incubation at extreme conditions
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Native protein
Molten globule
Unfolded protein
Aggregation
Amyloid fibril core structure
Further assembly of protofilaments and fibrils
FIGURE 13.15
Schematic representation of the amyloid formation process of proteins.
such as elevated temperature or acidic pH. Novel optical methods for the detection of amyloid fibril formation in proteins are of great importance with respect to the long-term stability and production of peptide pharmaceuticals in commercial pharmaceutical formulations used for the treatment of various diseases. Novel conformation-sensitive optical methods for the detection of formation of amyloid fibrils in bovine insulin (BI) and chicken lysozyme (CL) based on conformational changes of an anionic polythiophene derivative or a zwitterionic oligoelectrolyte were recently reported [28,61]. The technique is based on noncovalent assembly of the conjugated polyelectrolyte and the proteins (Figure 13.16). Depending on the conformation of the protein, different emission spectra from the conjugated polyelectrolyte are observed (Figure 13.17) [28,61]. The detection can also be observed by absorption and visual inspection (Figure 13.17), and this can be useful for the development of simple screening methods for the detection of amyloid fibrils. The conjugated polyelectrolytes are bound the native form of the proteins and to the amyloid fibrillar form of the proteins. These forms can easily be distinguished due to the conformational changes of the polyelectrolyte backbone upon binding to the different forms of the proteins, as minor perturbations of the geometry of the polyelectrolyte backbone can be reflected as alterations of the electronic structure of the conjugated backbone. Thus, binding of the polyelectrolyte to different forms of proteins will give rise to different optical features for the conjugated polyelectrolyte. This is an improvement to small dyes, used today, as these probes only change in optical feature whether they are free in solution or binding to pockets in the protein or to the surface of the protein. So far, it has been shown that conjugated polyelectrolytes can be used to distinguish between the native form of proteins and the amyloid fibrillar form of proteins. However, the use of conjugated polyelectrolyte might offer a novel approach to discriminate between different conformational structures observed during the amyloid formation processes.
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+
Native helical Twisted and separated protein polyelectrolyte chains Enhanced emission
Protein
Anionic conjugated polyelectrolyte
β−sheet containing Planar and aggregated polyelectrolyte chains Protein Decreased emission
FIGURE 13.16 Description of the detection of amyloid fibrils in proteins with an anionic conjugated polyelectrolyte. (From Nilsson, K.P.R., Herland, A., Hammarstrom, P., and Inganas, O., Biochemistry, 44, 3718, 2005.) 0.3
Absorption (m2)
0.25 0.25 0.15 0.1 0.05 0 340
440
390
490 540 Wavelength (nm)
590
640
1.80E+06 1.60E+06
Fluorescence (cps)
1.40E+06 1.20E+06 1.00E+06 8.00E+05 6.00E+05 4.00E+05 2.00E+05 0.00E+00 450
500
550
600
650
700
Wavelength(nm)
FIGURE 13.17 Absorption (top) and emission spectra (bottom) of PTAA (^), PTAA–native bovine insulin (&) and PTAA–amyloid fibrillar bovine insulin ().
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Amyloid insulin fibrils
Collagen
Collagen
Amyloid insulin fibrils
FIGURE 13.18 (See color insert following page 8-22.) Fluorescence images of a mixture of collagen and amyloid fibrillar bovine insulin stained with PONT (left) and PTAA (right).
The technique has so far been demonstrated for amyloid fibril formation in vitro, but initial experiments have shown that conjugated polyelectrolytes can be used as an amyloid-specific probe in histological staining of tissue samples (work in progress). As an example, a mixture of collagen and amyloid fibrillar bovine insulin was also stained by conjugated polyelectrolytes and the fluorescence images (Figure 13.18) clearly show distinct changes in color from the conjugated polyelectrolytes, depending onto which protein the conjugated polyelectrolyte is bound. This phenomenon is likely due to the intrinsic polyproline type II helical structure of collagen and the intrinsic b-sheet structure of the amyloid fibrils. In comparison to the polythiophene acetic acid (PTAA) fluorescence signal bound to amyloid fibrils, binding of PTAA to fibers of collagen showed a spectrum very similar to that of native bovine insulin [61]. The systems described in this chapter are mainly very simple chemical systems containing the desired molecules, mostly in an in vitro format. So the question remains: Can the sensory performance of conjugated polyelectrolytes be utilized in more complex systems, such as cells, blood samples, or tissue sections? We have just begun to explore the answer to this question and so far the results are quite promising. We have now demonstrated that staining of DNA in fibroblast preparations is feasible, and also that staining of misfolded proteins inside tissue sections is fully possible (work in progress). It is possible that these novel tools for biology may also lead to new procedures for the research and clinical laboratories.
References 1. Zhou, Q., and T.M. Swager. 1995. Methodology for enhancing the sensitivity of fluorescent chemosensors: Energy migration in conjugated polymers. J Am Chem Soc 117 (26):7017–7018. 2. Zhou, Q., and T.M. Swager. 1995. Fluorescent chemosensors based on energy migration in conjugated polymers: The molecular wire approach to increased sensitivity. J Am Chem Soc 117 (50):12593–12602. 3. Swager, T.M. 1998. The molecular wire approach to sensory signal amplification. Acc Chem Res 31 (5):201–207. 4. Chen, L.H., D.W. McBranch, H.L. Wang, R. Helgeson, F. Wudl, and D.G. Whitten. 1999. Highly sensitive biological and chemical sensors based on reversible fluorescence quenching in a conjugated polymer. Proc Natl Acad Sci USA 96 (22):12287–12292. 5. McQuade, D.T., A.E. Pullen, and T.M. Swager. 2000. Conjugated polymer-based chemical sensors. Chem Rev 100 (7):2537–2574. 6. Faid, K., and M. Leclerc. 1998. Responsive supramolecular polythiophene assemblies. J Am Chem Soc 120 (21):5274–5278.
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28. Herland, A., K.P.R. Nilsson, J.D.M. Olsson, P. Hammarstrom, P. Konradsson, and O. Inganas. 2005. Synthesis of a regioregular zwitterionic conjugated oligoelectrolyte, usable as an optical probe for detection of amyloid fibril formation at acidic pH. J Am Chem Soc 127 (7):2317–2323. 29. Nilsson, K.P.R., and O. Inganas. 2003. Chip and solution detection of DNA hybridization using a luminescent zwitterionic polythiophene derivative. Nature Mater 2 (6):419–424. 30. Stork, M., B.S. Gaylord, A.J. Heeger, and G.C. Bazan. 2002. Energy transfer in mixtures of water-soluble oligomers: Effect of charge, aggregation, and surfactant complexation. Adv Mater 14 (5):361–366. 31. Wang, J., D.L. Wang, E.K. Miller, D. Moses, G.C. Bazan, and A.J. Heeger. 2000. Photoluminescence of water-soluble conjugated polymers: Origin of enhanced quenching by charge transfer. Macromolecules 33 (14):5153–5158. 32. Wosnick, J.H., and T.M. Swager. 2000. Molecular photonic and electronic circuitry for ultrasensitive chemical sensors. Curr Opin Chem Biol 4 (6):715–720. 33. Harrison, B.S., M.B. Ramey, J.R. Reynolds, and K.S. Schanze. 2000. Amplified fluorescence quenching in a poly(p-phenylene)-based cationic polyelectrolyte. J Am Chem Soc 122 (35):8561–8562. 34. Wang, D.L., J. Wang, D. Moses, G.C. Bazan, and A.J. Heeger. 2001. Photoluminescence quenching of conjugated macromolecules by bipyridinium derivatives in aqueous media: Charge dependence. Langmuir 17 (4):1262–1266. 35. Jones, R.M., T.S. Bergstedt, D.W. McBranch, and D.G. Whitten. 2001. Tuning of superquenching in layered and mixed fluorescent polyelectrolytes. J Am Chem Soc 123 (27):6726–6727. 36. Fan, C.H., S. Wang, J.W. Hong, G.C. Bazan, K.W. Plaxco, and A.J. Heeger. 2003. Beyond superquenching: Hyper-efficient energy transfer from conjugated polymers to gold nanoparticles. Proc Natl Acad Sci USA 100 (11):6297–6301. 37. Kushon, S.A., K.D. Ley, K. Bradford, R.M. Jones, D. McBranch, and D. Whitten. 2002. Detection of DNA hybridization via fluorescent polymer superquenching. Langmuir 18 (20):7245–7249. 38. Heeger, P.S., and A.J. Heeger. 1999. Making sense of polymer-based biosensors. Proc Natl Acad Sci USA 96 (22):12219–12221. 39. Kushon, S.A., K. Bradford, V. Marin, C. Suhrada, B.A. Armitage, D. McBranch, and D. Whitten. 2003. Detection of single nucleotide mismatches via fluorescent polymer superquenching. Langmuir 19 (16):6456–6464. 40. Gaylord, B.S., A.J. Heeger, and G.C. Bazan. 2003. DNA hybridization detection with water-soluble conjugated polymers and chromophore-labeled single-stranded DNA. J Am Chem Soc 125 (4):896–900. 41. Liu, B., and G.C. Bazan. 2004. Interpolyelectrolyte complexes of conjugated copolymers and DNA: Platforms for multicolor biosensors. J Am Chem Soc 125:1942. 42. Xu, Q.H., B.S. Gaylord, S. Wang, G.C. Bazan, D. Moses, and A.J. Heeger. 2004. Time-resolved energy transfer in DNA sequence detection using water-soluble conjugated polymers: The role of electrostatic and hydrophobic interactions. Proc Natl Acad Sci USA 101 (32):11634–11639. 43. Wang, S., B.S. Gaylord, and G.C. Bazan. 2004. Fluorescein provides a resonance gate for FRET from conjugated polymers to DNA intercalated dyes. J Am Chem Soc 126 (17):5446–5451. 44. Liu, B., S. Baudrey, L. Jaeger, and G.C. Bazan. 2004. Characterization of TectoRNA assembly with cationic conjugated polymers. J Am Chem Soc 126 (13):4076–4077. 45. Gaylord, B.S., M.R. Massie, S.C. Feinstein, and G.C. Bazan. 2005. SNP detection using peptide nucleic acid probes and conjugated polymers: Applications in neurodegenerative disease identification. Proc Natl Acad Sci USA 102 (1):34–39. 46. Liu, B., and G.C. Bazan. 2005. Methods for strand-specific DNA detection with cationic conjugated polymers suitable for incorporation into DNA chips and microarrays. Proc Natl Acad Sci USA 102 (3):589–593. 47. Fan, C.H., K.W. Plaxco, and A.J. Heeger. 2003. High-efficiency fluorescence quenching of conjugated polymers by proteins. J Am Chem Soc 124:5642–5643. 48. Pinto, M.R., and K.S. Schanze. 2004. Amplified fluorescence sensing of protease activity with conjugated polyelectrolytes. Proc Natl Acad Sci USA 101 (20):7505–7510.
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49. Kumaraswamy, S., T. Bergstedt, X.B. Shi, F. Rininsland, S. Kushon, W.S. Xia, K. Ley, K. Achyuthan, D. McBranch, and D. Whitten. 2004. Fluorescent-conjugated polymer superquenching facilitates highly sensitive detection of proteases. Proc Natl Acad Sci USA 101 (20):7511–7515. 50. Dwight, S.J., B.S. Gaylord, J.W. Hong, and G.C. Bazan. 2004. Perturbation of fluorescence by nonspecific interactions between anionic poly(phenylenevinylene)s and proteins: Implications for biosensors. J Am Chem Soc 126 (51):16850–16859. 51. Rininsland, F., W.S. Xia, S. Wittenburg, X.B. Shi, C. Stankewicz, K. Achyuthan, D. McBranch, and D. Whitten. 2004. Metal ion-mediated polymer superquenching for highly sensitive detection of kinase and phosphatase activities. Proc Natl Acad Sci USA 101 (43):15295–15300. 52. Wosnick, J.H., C.M. Mello, and T.M. Swager. 2005. Synthesis and application of poly(phenylene ethynylene)s for bioconjugation: A conjugated polymer-based fluorogenic probe for proteases. J Am Chem Soc 127 (10):3400–3405. 53. Nilsson, K.P.R., J. Rydberg, L. Baltzer, and O. Inganas. 2004. Twisting macromolecular chains: Selfassembly of a chiral supermolecule from nonchiral polythiophene polyanions and random-coil synthetic peptides. Proc Natl Acad Sci USA 101 (31):11197–11202. 54. Leclerc, M. 1999. Optical and electrochemical transducers based on functionalized conjugated polymers. Adv Mater 11 (18):1491–1498. 55. Kumpumbu-Kalemba, L., and M. Leclerc. 2000. Electrochemical characterization of monolayers of a biotinylated polythiophene: Toward the development of polymeric biosensors. Chem Commun 19:1847–1848. 56. Ho, H.A., and M. Leclerc. 2004. Optical sensors based on hybrid aptamer=conjugated polymer complexes. J Am Chem Soc 126 (5):1384–1387. 57. Dore, K., S. Dubus, H.A. Ho, L. Levesque, M. Brunette, G. Corbeil, M. Boissinot, G. Boivin, M.G. Bergeron, D. Boudreau, and M. Leclerc. 2004. Fluorescent polymeric transducer for the rapid, simple, and specific detection of nucleic acids at the zeptomole level. J Am Chem Soc 126:4240–4244. 58. Bera-Aberem, M., H.A. Ho, and M. Leclerc. 2004. Functional polythiophenes as optical chemo- and biosensors. Tetrahedron 60 (49):11169–11173. 59. Nilsson, K.P.R., J. Rydberg, L. Baltzer, and O. Inganas. 2003. Self-assembly of synthetic peptides control conformation and optical properties of a zwitterionic polythiophene derivative. Proc Natl Acad Sci USA 100 (18):10170–10174. 60. Nilsson, K.P.R., and O. Inganas. 2004. Optical emission of a conjugated polyelectrolyte: Calciuminduced conformational changes in calmodulin and calmodulin–calcineurin interactions. Macromolecules 37 (24):9109–9113. 61. Nilsson, K.P.R., A. Herland, P. Hammarstrom, and O. Inganas. 2005. Conjugated polyelectrolytes: Conformation-sensitive optical probes for detection of arnyloid fibril formation. Biochemistry 44 (10):3718–3724. 62. Raymond, F.R., H.A. Ho, R. Peytavi, L. Bissonnette, M. Boissinot, F.J. Picard, M. Leclerc, and M.G. Bergeron. 2005. Detection of target DNA using fluorescent cationic polymer and peptide nucleic acid probes on solid support. BMC Biotechnol 5:10. 63. Wang, D.L., X. Gong, P.S. Heeger, F. Rininsland, G.C. Bazan, and A.J. Heeger. 2002. Biosensors from conjugated polyelectrolyte complexes. Proc Natl Acad Sci USA 99 (1):49–53. 64. Inganas, O., W.R. Salaneck, J.E. Osterholm, and J. Laakso. 1988. Thermochromic and solvatochromic effects in poly(3-hexylthiophene). Synth Met 22 (4):395–406. 65. Cheung, W.Y. 1980. Calmodulin plays a pivotal role in cellular regulation. Science 207:19–27. 66. Berggren, M., P. Bergman, J. Fagerstrom, O. Inganas, M. Andersson, H. Weman, M. Granstrom, S. Stafstrom, O. Wennerstrom, and T. Hjertberg. 1999. Controlling inter-chain and intra-chain excitations of a poly(thiophene) derivative in thin films. Chem Phys Lett 304 (1–2):84–90. 67. Kretsinger, R.H. 1980. Structure and evolution of calcium-modulated proteins. CRC Crit Rev Biochem 8:119–174. 68. Zhang, M., T. Tanaka, and M. Ikura. 1995. Calcium-induced conformational transition revealed by the solution structure of apo calmodulin. Nat Struct Biol 2:758–767.
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69. Kuboniwa, H., N. Tjandra, S. Grzesiek, H. Ren, C.B. Klee, and A. Bax. 1995. Solution structure of calcium-free calmodulin. Nat Struct Biol 2:768–776. 70. Carrell, R.W., and D.A. Lomas. 1997. Conformational disease. Lancet 350:134–138. 71. Prusiner, S.B. 1997. Prion diseases and the BSE crisis. Science 278:245–251. 72. Dobson, C.M. 1999. Protein misfolding, evolution and disease. Trends Biochem Sci 24:329–332. 73. Kelly, J.W. 2002. Toward an understanding of amyloidogenesis. Nat Struct Biol 9:323–325.
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14 Conjugated Polymers for Microelectromechanical and Other Microdevices 14.1
Introduction..................................................................... 14-1 What Are MEMS and Microsystems? . Current MEMS Applications . Conventional MEMS Materials . MEMS Fabrication Techniques
14.2
Conjugated Polymers for MEMS ................................... 14-6 Properties of Conjugated Polymers at the Microscale and Nanoscale . Processing and Fabrication Techniques for Conjugated Polymer Microdevices
14.3
Chemical and Biological Sensors ................................. 14-13 Current MEMS Chemical Sensor Principles Selection . Conjugated Polymer Sensors
14.4
.
Materials
Actuators ........................................................................ 14-17 Current MEMS Actuator Principles . Materials Selection . Conjugated Polymer Microactuators
14.5
Microfluidic Systems..................................................... 14-21 Current Microfluidics: Principles and Methods . Materials Selection . Conjugated Polymer Pumps and Valves
14.6
Geoffrey M. Spinks and Elisabeth Smela
14.1
Energy Storage Systems ................................................ 14-23 Current Microfabricated Energy Storage Methods . Materials Selection . Conjugated Polymer Batteries and Capacitors
14.7
Future Challenges and Possibilities.............................. 14-25
Introduction
14.1.1 What Are MEMS and Microsystems? The acronym MEMS stands for microelectromechanical systems, which have lateral dimensions in the range of 1–1000 mm. MEMS are also sometimes referred to as microsystems technology (MST) and the fabrication techniques as micromachining. MEMS devices are fabricated using techniques that were originally developed for the microelectronics industry but that have been adapted and extended to
14-1
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construct a wide variety of devices that accomplish a large number of engineering functions. In particular, MEMS devices are used either as sensors or as actuators, and sometimes these elements are coupled together to produce a microsystem. MEMS may or may not contain moving parts. For reviews of the MEMS field, see for example Refs. [1–3]. Perhaps the most ubiquitous example of such systems is the air bag deployment mechanisms incorporated into most modern automobiles. These systems use microaccelerometers to trigger the air bag release at the appropriate time. Another common MEMS device is the printhead used in ink-jet printers, which includes the pump system and the nozzles. Miniaturized pressure sensors serve as an example to illustrate the advantages of MEMS devices. Very small MEMS pressure sensors have replaced metal diaphragms with bonded strain gauges, allowing them to be widely used in automobiles (e.g., fuel tanks, fuel injectors, engine oil [4], and soon tires) to optimize vehicle operation and to inform the driver about the operating conditions. The smaller size, higher reliability, and lower cost were important in these applications so that the sensors could be incorporated into components throughout the vehicle and operate throughout the life of the vehicle. MEMS pressure sensors utilize a diaphragm only a few micrometers thick whose deflection depends on the surrounding fluid pressure, and they incorporate circuitry that determine the deflection based on readings from piezoresistors implanted in the diaphragm [5]. As most silicon-based fabrication steps are now routine processes, it is possible to construct the mechanical and electrical components on the same chip. A complete device design is shown in Figure 14.1, and the fully packaged sensor element is smaller than 1 cm. The MEMS industry now represents a multibillion dollar market in annual worldwide sales. Estimates of the value of the industry vary widely, however, mainly because of the blurry definition of a MEMS device and because of the difficulty of valuing the MEMS component of a larger system that is enabled by this technology. In some analyses, MEMS are considered to be devices containing moving parts in the micrometer to millimeter size range. Other reviews (including this chapter) consider MEMS to be indistinguishable from microsystems, which do not have moving parts (such as microfluidic channels). Another difficulty in estimating the size of the microsystems market is the fragmented nature of the industry, because MEMS devices have found their way into a plethora of products from scientific
Metal casing Piezoresistors (4) Wheatstone bridge
Wire bond
Metal pad
Silicon diaphragm
Dielectric layer Constraining base (Pyrex glass)
Interconnect
Die attach
Inlet for pressurized fluid
FIGURE 14.1 Illustration of a MEMS pressure sensor, including sensing element and associated transduction electronics. (From Hsu, T.-R. MEMS & Microsystems: Design and Manufacture, 1st ed., McGraw-Hill, Boston, 2002.)
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Signal transduction and processing unit
Information processing Sense
Actuator
Sensor
Intelligent material
Actuate
Energy storage / conversion
(a)
Power supply
(b)
FIGURE 14.2 (a) Components of a microsystem. (From Hsu, T.-R. MEMS & Microsystems: Design and Manufacture, 1st ed., McGraw-Hill, Boston, 2002.). (b) A smart material system (From Wallace, G.G., Spinks, G.M., KaneMaguire, L.A.P., and Teasdale, P.R., Conductive Electroactive Polymers: Intelligent Materials Systems, 2nd ed., CRC Press, Boca Raton, 2002.)
instruments to cell phones. Estimates from 1999 of the MEMS market put the value at several billion dollars [5,6], with over 80% of the market in silicon-based microsensors [5]. By 2005, the market was expected to reach $5.7 billion [7]. In contrast, the European organization Network of Excellence in Multifunctional Microsystems (NEXUS) estimated a market of $30 billion in 2000 for the broader category of microdevices, rising up to $68 billion in 2005 [7]. Some analysts warn against believing market projections for MEMS products [8], because major new products (including air bag sensors) have rapidly emerged. (that were not predicted by prior market analyses). Clearly, one important advantage of producing a microdevice is miniaturization. The air bag accelerometers are again an illustrative example. However, the biggest advantage is improved performance at lower cost. Each MEMS accelerometer is no larger than the period at the end of this sentence [6] and in 1999 cost less than $10. In contrast, the prior motion sensors consisted of five components, which are several centimeters in diameter, and they cost $18 each. Furthermore, the MEMS accelerometer includes self-test capability to ensure that the device is in working order, which the prior devices did not. Just as ICs have become more powerful, ubiquitous, and cheaper, it is also expected that the application areas for MEMS will steadily increase while the cost of MEMS devices will continue to decline. In fact, much of the hype surrounding the MEMS industry is due to perceived parallels with the integrated circuit industry. The development of microsystems has uncanny parallels to the rationale used to promote conjugated polymer (CP) technology. Conjugated polymers are synonymous with smart materials and smart systems. So too are MEMS and microdevices. Hsu [5] defines a microsystem as shown in Figure 14.2a with sensor and actuator elements linked by a signal transduction and processing unit. Virtually identical concepts have been highlighted as the basis of smart material systems (Figure 14.2b). In both systems the fundamental idea is the same: the sensor detects a change in the local environment (such as temperature, pressure, chemical species), and the actuator produces an output response (such as a movement or change in color). As noted by Karen Markus (Cronos Integrated Microsystems, NC), ‘‘Computers think and think and think. But MEMS are becoming the eyes, ears, nose, mouth, hands, and feet of computers.’’ [6]. Perhaps the best smart systems will be those microdevices that incorporate smart materials, such as conjugated polymers. This chapter explores the tentative steps already taken toward that goal.
14.1.2 Current MEMS Applications The development of MEMS devices can be traced back to the 1970s, when micromachining of silicon was first investigated for the manufacture of mechanical devices. The MEMS market is now dominated
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Automotive industry
Health care products
Aerospace industry
Industrial products Consumer products
Safety systems: air bag deployment, antilock braking, suspension, navigation Engine and power train: manifold control, airflow control, fuel injection, transmission pressure Comfort: seat control, temperature and humidity, security, defogging controls, navigation Vehicle diagnostics: coolant temperature, oil pressure, tire pressure, transmission fluid, fuel level Disposable blood pressure transducers Angioplasty pressure sensors Infusion pump pressure sensors Catheter tip pressure sensors Medical process monitoring Kidney dialysis equipment Cockpit instrumentation: pressure sensors, airspeed, altimeters Microgyroscopes Sensors for fuel efficiency and safety Sensors for hydraulics, paint spray, agricultural sprays, refrigeration systems, air conditioning, water level Scuba diving watches and computers Bicycle computers Smart vacuum cleaners, washing machines, and other home appliances Smart toys
Source: From Hsu, T.-R. MEMS & Microsystems: Design and Manufacture, 1st ed., McGraw-Hill, Boston, 2002.
by four large applications: microfluidics, accelerometers, optical MEMS, and pressure sensors. Other types of sensors, RF MEMS, and actuators make up smaller markets. The end products that use microdevices are diverse, but a number of examples in various market categories are given in Table 14.1. The list of applications shows that the market is dominated by sensor applications, which are mainly variants of the pressure sensors and accelerometers developed some two decades ago. Emerging applications are also quite varied, but can be classified into one of three areas: chemical and biological sensors (including biochips and DNA chips [10]); microactuators; and power MEMS. Each of these three areas is described below, together with the benefits that conjugated polymers can potentially offer, including performance enhancements and increased functionality. It is likely that the first commercial adaptation of conjugated polymers to MEMS will occur through an evolutionary process involving existing devices.
14.1.3 Conventional MEMS Materials Given the birth of MEMS from the IC industry, the dominant material used in the early devices was silicon. The use of silicon as a substrate and structural material, and the use of polysilicon as a thin film structural material, has continued to the present day for several reasons. The microfabrication techniques for silicon are highly developed and flexible, the microfabrication equipment has been designed for silicon, the properties of silicon are very well known and can be tightly controlled, silicon has excellent mechanical properties, and silicon has an insulating native oxide that can be used as a sacrificial layer. Other thin film materials that are commonly used in MEMS include silicon nitride, metals, and conventional polymers, such as polyimide. Other substrates that are used for MEMS include quartz, glass, and polymers. Quartz is attractive primarily because it is piezoelectric, so that it may be used as a sensor and an actuator. Quartz is, however, more difficult to micromachine than silicon. Pyrex glass is used in conjunction with silicon wafers, primarily for packaging, because of its optical transparency and close match in coefficient of thermal expansion (CTE). Polymers such as polycarbonate have been adopted as substrates for microfluidics because of their low cost: channels require large areas, which makes silicon too expensive, and channels can be fabricated in polymers inexpensively by hot embossing. The advantages offered by silicon substrates are not required because microfluidic devices typically do not have complex
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microstructuring or moving parts. Polydimethylsiloxane (PDMS) is another polymer that is increasingly finding its way into MEMS, particularly for lab-on-a-chip applications, because of the ease of structuring it by molding and because of its biocompatibility. In all application areas, a key material requirement is processability. Fabrication requirements can dominate in the choice of materials for MEMS. This restriction has kept shape memory alloys (SMAs), for example, from finding widespread application, although SMA MEMS devices have been demonstrated in the laboratory.
14.1.4 MEMS Fabrication Techniques There are two main parts for the production of completed microsystems: microfabrication to produce the devices or systems, and packaging to protect the components and provide the necessary signal inputs and outputs. MEMS packaging is often more challenging than IC packaging, and is highly dependent on the individual application, and must be an integral part of the design process from the beginning. However, a discussion of packaging issues is outside the scope of this chapter (for a review of this topic refer to Madou [11]). The microfabrication methods for conjugated polymers (see Section 14.2) must be compatible with those for current MEMS devices in order for these materials to be adopted. It is, therefore, necessary to understand these processes, which are briefly described here. Microfabrication techniques for MEMS can be classified into two categories: bulk micromachining and surface micromachining (including Lithografie Galvanoformung Abformung [LIGA], which translates to lithography, electroplating, and molding). Bulk micromachining involves forming MEMS structures from the substrate itself through the removal of material to produce holes, channels, cantilevers, diaphragms, and other two- or threedimensional structures. In surface micromachining the substrate serves as a platform, and the devices are fabricated layer by layer from thin films. The use of a ‘‘sacrificial’’ layer allows structures to be partially freed from the substrate, allowing them to move. Structural layers are built-up over the sacrificial layer, which is later removed to produce, for example, suspended beams or cantilevers. Materials are deposited onto the surface of the substrate using a variety of techniques, from spin-coating of polymers to chemical vapor deposition of dielectrics. Metals are usually deposited by thermal evaporation, e-beam evaporation, or sputtering. Silicon dioxide can be produced by reacting the silicon itself at high temperatures in the presence of oxygen or water. Typical film thicknesses are tens of nanometers to 2 mm. In both surface and bulk micromachining, the most common method to pattern materials is selective etching (lift-off is another). Control over the spatial location of the etch is achieved through masking, and selectivity to the material to be etched is achieved through the choice of etchant. Masking is performed by photolithography, as shown in Figure 14.3. Photolithography is akin to photography, with the mask playing the role of the negative, photoresist the role of the silver emulsion, and the wafer the role of the paper. A layer of photoresist is spin-coated onto the substrate and exposed to UV light through a mask. The resist is then developed, producing a copy of the mask pattern in the resist (which can be positive or negative, as a slide or photo negative). The resist protects underlying material during the subsequent etch, and then it is removed. Etching is performed either by using wet chemicals or by using reactive gases (dry etching). The former may involve electrochemical reactions. The latter can involve chemical reactions, physical bombardment (sputtering), or a combination of the two. Both wet and dry etching can be either isotropic (the same in all directions, producing rounded features and mask undercutting) or anisotropic (directional). For example, wet etch rates for different crystal planes in crystalline materials can vary considerably. Anisotropic etches allow the formation of structures with high aspect ratio (height and width). The LIGA process is a method to produce high aspect ratio structures by adding material onto the surface rather than by subtracting it through etching. This is done by plating a metal onto a high aspect ratio mold produced by x-ray lithography (for a good review of this technology, see Chapter 6 of
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Photography Dark room
Photolithography Clean room
Visible light
UV light
Photo paper
Resist-covered layer Develop
Positive resist positive resist is broken down by UV light where it is exposed, it goes away upon development
FIGURE 14.3
Negative resist negative resist is cross-linked by UV light where it is exposed, it stays upon development
General procedure for photolithography.
Ref. [11]). The process begins with an electrically conductive layer on the substrate (or a substrate that is conducting). The substrate is coated with a thick layer of a special polymer resist, which is usually poly(methyl methacrylate) (PMMA), because of its favorable optical absorption and processing characteristics. The polymer is exposed through specially designed and fabricated masks. X-ray is used because it can penetrate deeply into the polymer and thereby produce high aspect ratio features. The polymer is then developed to produce the mold. The metal (e.g., nickel) is either electrolessly plated or electroplated onto the conducting layer to fill the mold. Complex three-dimensional structures can be built-up through multiple layers, and the final structure is freed by etching away the remaining resist. The metal structures can in turn be used as injection molds for the production of polymer MEMS. Because of the high cost and complexity of LIGA, variations of the technique have been developed for use in conventional clean rooms based on thick resist and UV lithography.
14.2
Conjugated Polymers for MEMS
Before considering the use of conjugated polymers for MEMS and microsystems, two things should be considered: the properties of these polymers at the microscale and the processing techniques that can be used to pattern them.
14.2.1 Properties of Conjugated Polymers at the Microscale and Nanoscale The general properties of conjugated polymers are exhaustively covered in other chapters of this handbook (see Chapter 7, Chapter 10, and Chapter 13). Some of the properties of conjugated polymers at the microscale are summarized here and in other works [12]. Two important properties relevant to the function of conjugated polymer devices are their conductivity and electroactivity. The conductivity is important to sensor devices because changes in conductivity produce the output signal. Although high conductivity is important, reducing resistive losses is the key requirement. For example, conjugated polymer batteries require efficient charge transport into and out of the polymer, and resistive losses within the polymer can reduce device performance. Batteries, actuators, and some sensors made from conjugated polymers operate through a redox cycle in which the efficiency of the electroactive processes is a key concern. Efficiency can be defined as the proportion of charge injected or removed from the polymer compared with the total charge available through
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complete charge and discharge. Often the efficiency of microscale redox systems is improved because the surface area to volume ratio is high, giving a large interface between the electrode and the electrolyte. The conductivity of conjugated polymers varies throughout the material, with highly conductive domains separated by less conductive material [13,14]. The highly conductive domains are typically tens of nanometers in size [15] and consist of more highly ordered, or even crystalline, regions. One attractive feature of miniaturization is the possibility of having the entire structure to consist of highly ordered material. Several reports have shown that the conductivity of CP nanostructures is higher by 1 to 2 orders of magnitude than that of corresponding macroscale films. Delvaux and coworkers [16], for example, found that template-synthesized [17] polyaniline micro- and nanotubules showed an enhanced conductivity when the tubule diameter was less than 200 nm. Several earlier studies had also shown the same results for polypyrrole (PPy) and polythiophenes [18,19]. In a recent study of the electronic structure of PPy nanotubes [20], it was found that the average conjugation length was larger, and the energy of the pp* transition is lower, in the nanotubes compared with films produced under identical conditions. In redox systems, microelectrodes are known to show certain advantages [21,22]. The double layer capacitance and the RC time constant are smaller with microelectrodes, so fast voltammetric measurements can be made. In addition, the depletion region around the electrode is small, so the current is not affected by movement of the solution, which means that microelectrodes can operate successfully in flowing streams, such as within microfluidic systems. Microelectrodes can also operate in resistive electrolytes (at low ion concentrations or low temperature) because the ohmic resistance of a cell is proportional to the electrode size. These advantages have been utilized in conducting polymer microelectrodes for the detection of chemical species [23,24], with one study [23] showing that a microelectrode had a lower detection limit than a corresponding macroelectrode. Additionally, the microelectrode could operate successfully in a low buffer strength environment (resistive), which was more typical of real test environments. In general, the electroactivity of conjugated polymers at the microscale has been found to be more efficient than at the macroscale. Transport of ions within the polymer is often the rate-limiting process during electrochemical reactions. Ding et al. [25] have studied the efficiency of polypyrrole electroactivity for different polymer film thicknesses deposited on platinum. They found that using short pulses, only a fraction (<20%) of the available charge was recovered from the polymer when the film thickness exceeded 1 mm. As much as 60% of the available charge could be transferred in PPy films, which is 70 nm thick. These results reflect the kinetic limitations occurring in conjugated polymers, and highlight the faster responses that can be achieved in microscale devices. Fast electrochemical responses have also been observed in individual polyaniline nanowires [26]. When the nanofiber length is reduced so that the fiber contains only one highly conducting domain, electrochemical switching between oxidized (conducting) and reduced (insulating) states was found to occur in 0.01 ms and at a single electrochemical potential. In longer nanofibers (and in bulk materials) the electrochemical switching occurs over a range of potentials due to both distribution of conjugation length and the convolution of sweep rate with the speed of ion transport.
14.2.2 Processing and Fabrication Techniques for Conjugated Polymer Microdevices The ability to fabricate micron-sized features from conducting polymers is a key requirement in their adoption in MEMS and microsystems. Conducting polymers can be deposited onto a substrate by electrodeposition or spin coating, and a number of techniques have been explored for micropatterning them. This section reviews the current state of microfabrication techniques for conducting polymers. Whereas a wide variety of patterning methods are available, not all are suitable for each conjugated polymer. Care needs to be taken during processing to limit the temperatures to which the conjugated polymer is exposed, as well as the solvents, etchants, and other chemical species.
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14.2.2.1 Deposition
14.2.2.1.1
Spin-Coating
Depositing thin, uniform polymer films by spin-coating is a widely utilized technique in microfabrication. Resist layers 1–2 mm thick are typically deposited by spin-coating on substrates such as silicon wafers, although films less than 0.1 mm can also be made [11]. The quality of films produced by spincoating depends on the solution viscosity and the solvent volatility. Spin-coating is used for the preparation of films and coatings of soluble conducting polymers such as polyaniline and polythiophenes. Some applications of spin-coating in the preparation of conducting polymer devices are described in recent publications [27–29].
14.2.2.1.2 Electrodeposition Electrochemical processes are widely used in MEMS fabrication for selective etching, the formation of porous silicon, and electroplating of metals. Electrodeposition is also a widely used technique for preparing conjugated polymers, either as electrode coatings or as freestanding films (after removal from the electrode) because many CPs are not soluble, and none are melt-processable. Chemical oxidation can also be adapted to produce films by coating the surface with oxidant and exposing the substrate to the monomer vapor (or sometimes the reverse procedure is used) [30,31]. Although chemical oxidation is a simple technique, electrodeposition has several advantages, including control of the dopant type (and, therefore, subsequent properties), greater control over film morphology, and ability to produce thicker coatings (tens of microns). Although chemical oxidation usually produces powders, this technique can be adapted to produce patterned coatings by first patterning the oxidant on the substrate surface. Electrodeposition requires an inert electrode that will not oxidize under the polymerization conditions. Usually, gold, platinum, carbon, or stainless steel is used. Electrodeposition of CPs begins at the electrodes, but, depending on the surface, can spread laterally to generate conducting paths between adjacent electrodes [32]. The electrodeposition process requires an electrochemical cell: a working electrode (onto which the polymer is deposited), an auxiliary electrode (to provide the current), a reference electrode (if voltage-control is used), and an electrolyte (a salt-containing solvent). Electrodes and interconnects must, therefore, be included to produce a CP MEMS device, as well as to cycle it (if it is a device such as CP actuator). Electrochemical deposition of conjugated polymers occurs easily, but the final properties are sensitive to the deposition conditions, including supporting electrolyte, temperature, electrical stimulus (galvanostatic, potentiostatic, or potentiodynamic), and electrode material. MEMS developers need to be familiar with how deposition conditions affect film properties, such as conductivity, Young’s modulus, and speed. Further details on the effects of electropolymerization conditions on the structure and properties of conducting polymers can be found elsewhere [9]. 14.2.2.2 Patterning There are several ways to pattern CPs, as reviewed previously [33–35]. The three main methods are as follows: 1. Patterning the electrode onto which the CP is deposited; 2. Depositing the CP everywhere and then etching it; 3. Printing the polymer. The first and second methods are used with electrodeposited films, and the second and third with soluble CPs. These and other methods are described in the following sections.
14.2.2.2.1 Selective Electrodeposition By polymerizing the conjugated polymer in selected places on the surface, deposition and patterning are combined into a single step. There are two main methods for this: deposition onto patterned electrodes
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and growth in photoresist molds. Both of these are batch-fabrication, parallel process that are easy to implement and are already used to electroplate metals. One way to pattern a conjugated polymer is to pattern the metal electrode onto which it deposits, and Figure 14.4a illustrates the method. The metal is patterned, for example, photolithographically by the application, exposure, and development of photoresist (Figure 14.4a1) followed by metal etching (Figure 14.4a2) and removal of the resist (Figure 14.4a3), and then the polymer is electrochemically deposited (Figure 14.4a4). This is one of the simplest methods. The main drawback is that the film thickness is not uniform: higher electric fields and availability of reactants at the edges of the electrodes produce thicker deposits there. Also, the differential adhesion technique, which is an alternative to the sacrificial layer, cannot be used. If the surface is modified to be hydrophobic [33,36], then lateral growth can be substantially faster than vertical growth, and adjacent electrodes can be connected (Figure 14.4a5). Another method is template polymerization, for example, using photoresist as a mold, as illustrated in Figure 14.4b1. The conjugated polymer will only deposit materials in the openings (Figure 14.4a2), and the resist can subsequently be removed (Figure 14.4a3) with ethanol or by blanket exposure and development. A limitation is that films are restricted to approximately the thickness of the photoresist. It should be noted that, in general, deposition in templates results in different properties at the template walls than on the electrode, particularly if the template is a polymer or some other material that interacts with the CP. Growth rate of walls is often faster, although this depends on surface energy (as does the lateral growth rate, Figure 14.4a4 and Figure 14.5). With this method, the edges have been found not to actuate with the same strain as the center [37]. Patterning of the electrode can subsequently be performed using the CP as a mask. If the metal etchant is damaging to the CP, a resist layer can be used to protect it, although this can result in small misalignments (Figure 14.4b4). Other templates have also been used. Polypyrrole nanowires have been produced by growing the polymer in porous alumina. Highly porous conducting polymers have also been produced using the inverse opal method [38], in which the polymer is deposited around a matrix of tightly packed spheres, which form the synthetic opal. When the spheres are removed, a highly porous film is left with
(a)
(b)
(c)
(d)
(1)
(2)
(3)
(4)
(5)
FIGURE 14.4 Most frequently used patterning methods for conjugated polymer films. (a) Deposition on patterned electrodes. (b) Deposition within a template. (c) Etching. (d) Printing. For details, refer to the text. Vertical dimensions and other features have been exaggerated for illustrative purposes.
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Conjugated Polymers: Processing and Applications
Commercially available metal-coated plastic substrate
Print dielectric layer (polyimide)
Print polymer semiconductor (polythiophene) Print drain and source electrodes (conducting ink)
(b)
i E
Polymer film 2 mm
Supporting substrate 5 mm
FIGURE 14.5 Examples of conducting polymer devices produced by screen printing: (a) procedure for preparing a polythiophene-based field effect transistor. (From Bao, Z., Rogers, J.A., and Katz, H.E., J. Mater. Chem., 9, 1895, 1999. With permission.); (b) polypyrrole-based chemical sensor shown schematically (left) and with photographs of screen-printed carbon electrodes (top right) and with polypyrrole subsequently deposited over such electrodes (bottom right). (From Shepherd, R.L., Barisci, J.N., Collier, W.A., Hart, A.L., Partridge, A.C., and Wallace, G.G., Electroanalysis, 14, 575, 2002. With permission.)
controlled pore size and spacing. Similar techniques have been proposed for the formation of conducting polymer interconnected with wires in three-dimensional chips [39] and for polythiophene electrochromic displays [40]. A variation of the technique uses self-assembled monolayers (SAMs) to mask the electrode. In one example, alkane thiols were applied to a gold surface using microcontact printing. Subsequent electropolymerization of pyrrole [41] or aniline [42] produced patterns 10–100 mm in width. Resolution of 2 mm can be achieved for films 10 nm thick [43], but thicker films might undergo lateral growth. Patterning of SAMs has also been achieved by irradiation through a mask, and this has been used to selectively electrodeposit polypyrrole, polyaniline, and poly(3-methylthiophene) [44,45]. Patterns made from polymer brushes grown on the substrate have also been used as masks [46]. A third category of patterned electrodeposition uses a scanning counter electrode moved over the working electrode surface. Polypyrrole [47] and polyaniline [48] tracks of several millimeters in length and tens of microns in width have been produced using this method. Although complex patterns can be prepared, this technique is a serial process and, therefore, inherently slow, so it will likely find only limited applications, such as for prototyping [49].
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14.2.2.2.2 Photolithography and Etching The most standard way to pattern conjugated polymer films is to deposit the film over the entire surface (Figure 14.4c1), mask it (Figure 14.4c2), and then etch it by reactive ion etching (RIE) (Figure 14.4c3), usually in an oxygen plasma. This method produces patterned films of uniform thickness and line widths that are limited essentially by the optics of the mask and the mask aligner. Using the resist and CP as a mask for subsequent etching of the metal (Figure 14.4c4) gives perfect alignment (this is a so-called ‘‘self-aligned’’ process). Submicron resolution of thin films is possible, because RIE is directional (anisotropic). This method works very well for most conjugated polymers, including polypyrrole and polyaniline. An oxygen plasma will not, however, etch polyethylenedioxythiophene (PEDOT). Regular photoresist can be used as a mask, but its etch rate is comparable to that of the polymer. In order to etch thicker conjugated polymer films, another masking layer must be used, such as a gold film. (Gold etchant does not swell polypyrrole, but aluminum etchant does, causing the CP to delaminate if it is exposed to this solution. Therefore, aluminum should not be used as a mask.) Other photolithographic techniques have also been developed for micropatterning of conducting polymers. For example, radiation-induced cross-linking has been used in the patterning of polyaniline [50] and its derivatives [51], as well as poly(3-alkyl thiophenes) [52]. Photobleaching and photodoping have also been demonstrated for patterning. The former does not physically remove the polymer, but renders it insulating and electrochemically inactive. The end result is regions of different functionality within the polymer film [50,53]. Photodoping, through the incorporation of photoacid generator (PAG) and exposure to UV light, has been used to protonate polyaniline to form submicron patterns [54]. Renak et al. [55] have used this technique to prepare an array of photoluminescent spots from a poly(phenylene vinylene) (PPV) copolymer. 14.2.2.2.3
Printing
Soluble polymers are amenable to printing (Figure 14.4d). Printing is an additive process in which material is deposited only where it is needed; the deposition and patterning steps are thus, combined. However, alignment is substantially less accurate than with photolithography, and, therefore, printing is limited to single layers or relatively large devices. Printing can be applied to a variety of substrates with high throughput. For the above-mentioned reasons, printing of soluble conducting polymers is seen as a key processing technology in the development of polymer-based electronics. Clearly, these advantages can also be applied to conducting polymer MEMS. The ancient technique of screen printing has been applied for the production of conducting polymer devices [56]. Soluble ink is squeezed through a screen upon which the desired pattern has been formed and then onto a suitable substrate. Evaporation of the solvent leaves the conducting polymer pattern. Examples of demonstration devices produced by screen printing of conducting polymers are shown in Figure 14.5. Screen printing is rapid once the screen has been constructed, but resolution below 10 mm is difficult. Also, inks need a reasonably high viscosity to prevent spreading. The low solubility of many conducting polymers means that solution viscosities can be too low for successful screen printing. Ink-jet printing is a more recent development that produces higher resolution in the size and shape of features than screen printing and has also been adopted for patterning of conducting polymers [58,59]. Polypyrrole films for chemical sensing were printed this way in 1998 [60]. The resolution achievable in ink-jet printing depends on the ejected droplet volume and the degree of spreading of the droplet on the substrate. The latter is determined by wetting, binding, and evaporation characteristics of the ink. The droplet volume is influenced by the ink viscosity. Early ink-jet printing of conducting polymers used commercial ink-jet printers and was unable to achieve small feature sizes because the inks and substrates were not optimized for this process. Developments in the design of ink-jet printing heads [61] and better understanding of drop formation using conducting polymer inks [62] has led to the production of efficient ink-jet printing systems for conducting polymer light-emitting diodes (LEDs) [63]. The first use of ink-jet printing to construct MEMS has recently been described by Fuller et al. [64]. Metal nanoparticles suspended in a solvent were deposited, followed by sintering at elevated temperatures.
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Printing multiple layers could even produce three-dimensional structures. A thermal actuator has been fabricated with this technique. MEMS fabrication routinely makes use of a process known as microcontact printing, pioneered by Whitesides and coworkers, which can achieve microscale resolution. A stamp is constructed from micromolded PDMS, and the material to be patterned is inked onto the stamp surface. When the stamp is brought into contact with the substrate, only the material on the high points of the stamp are transferred. The method has been used to prepare lines of light-emitting polythiophenes with 100 mm width and thickness in the micron range [65]. Microcontact printing has also been used to mask substrates so that subsequent deposition of conducting polymer only occurs on certain areas [66–69]. In one recent example [70], PPy patterns were prepared on a silicon substrate with line widths of 25 mm, heights of 2 mm, and lengths in excess of 2 mm. Conductivities of 0.5 S=cm were measured for these features. In addition, it has been demonstrated that microcontact printing can be used to chemically modify a conducting polymer so that printed areas become insoluble. The untouched areas can be dissolved away to leave a conducting pattern [71]. In yet another variation, the printed ink was specifically chosen to contain reactive groups that could be covalently bonded to the conducting polymer. Using this method, an amino-functionalized PPV was patterned on a gold surface that had previously been printed with 5 5 mm2 of 16-mercaptohexandecanoic acid. The patterned surface was immersed in a solution of the polymer, followed by rinsing and drying to give the patterned polymer shown in Figure 14.6 [72]. Finally, a mold (usually PDMS) reversibly adhered to a substrate produces channels into which a polymer solution can be introduced using capillary action. This technique has been used to manufacture polyaniline field effect transistors [73]. Nanoscale patterning can be achieved by using dip pen nanolithography (DPN) [74]. In this method, the imaging tip of an atomic force microscope (AFM) is coated in the material to be deposited. The tip is brought into contact with the substrate and moved so as to transfer the coating material to the substrate. The transfer process requires the coating material to preferentially adhere to the substrate rather than the tip. The quality of the patterns depend on the environment (temperature, humidity), tip coating procedures, tip writing speeds, and ink–substrate interactions [74,75]. The direct application of conducting polymers using DPN has utilized water-soluble polyaniline and polypyrrole applied to silicon surfaces by electrostatic interactions [75]. In another variation [76], a reactive ink of pyrrole monomer with perchloric acid diluted in tetrahydrofuran was used to prepare conducting polypyrrole tracks on a silicon surface. 30 nm 68 60 66
15 µm
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FIGURE 14.6 AFM image (left) and height profile (right) between triangle markers shown in the image of PPV deposited on a gold surface. Before PPV deposition the gold surface was patterned using microcontact printing so that the PPV deposited in selected areas. The square features shown in the image are regions with no PPV. The height profile shows the PPV layer to be approximately 6–7 nm thick. (From Liang, Z.Q., Rackaitis, M., Li, K., Manias, E., and Wang, Q., Chem. Mater., 15, 2699, 2003. With permission.)
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Many of the printing techniques described above use polymer solutions as ink, and it is recognized that solvent incompatibilities can mean that such techniques are not always suitable. A dry printing technique using thermal imaging has been used to construct an array of thin film transistors with features as small as 15 mm [77]. The thermal imaging technique involves the heating of selected areas of the material to be patterned using an infrared laser. The heated areas are transferred by contact to a nearby substrate. Multiple layers of materials can be built-up using the thermal imaging technique without the complication of solvents dissolving previously deposited layers.
14.2.2.2.4
Others
Although not a bona fide patterning method, phase separation of block copolymers has been used to form templates for selective deposition. Phase separation occurs at micrometer to nanometer length scales. Polymer blocks with different surface energies can be used to allow and prevent deposition of conducting polymers. In one example [78], polypyrrole was chemically polymerized and deposited on a surface previously coated with a block copolymer of styrene, vinyl pyridine, and surfactants. The polymer is deposited preferentially on one of the copolymer domains to give PPy regions with lateral dimensions in the range of 100 nm. Domains with different shapes were demonstrated, but the technique cannot be used to give arbitrary circuit or device structures. Innovative microstructures have been produced by self-assembly during electropolymerization in the presence of surfactants. Dai and coworkers have recently shown that bowl-shaped ‘‘microcontainers’’ of polypyrrole can be produced by stabilizing H2 gas bubbles on the electrode surface [79]. Carboxylic acid dopants have also been used, resulting in hollow nanotubes of polyaniline [80]. In the remainder of this chapter, we will discuss how the unique properties and functions of conjugated polymers can be incorporated into various types of microsystems, starting with sensors. The limitations of current MEMS materials are discussed in the context of specific examples. A more speculative discussion of the possible use of conjugated polymers in entirely new microsystems is considered at the end of the chapter.
14.3
Chemical and Biological Sensors
One of the areas in which conjugated polymers may improve on the performance of current MEMS materials would be chemical and biological sensing. Miniaturized sensors are expected to find applications in a wide variety of areas. Sensors for chemical and biological species are particularly important in environmental monitoring, the food industry, health care, and border security. The availability of miniaturized and inexpensive chemical and biological sensors would enable these systems to be widely distributed and extensively used. The continuous monitoring of pollution in waterways is one example where ubiquitous chemical sensors would have a major impact. Networks of chemical sensors are envisaged for continuous, autonomous monitoring of chem-bio hazards and pollutants [81,82]. Miniature, low-cost microsensors are developed for incorporation into the packaging or shipping containers for food and pharmaceuticals to indicate spoilage or tampering. Civil engineering applications include the early detection of corrosion in structures, such as bridges and aircraft [83,84]. Miniaturization will also have an impact in health care with the development of lab-on-a-chip systems for rapid diagnosis.
14.3.1 Current MEMS Chemical Sensor Principles To consider the possibilities of using conjugated polymers in MEMS sensors, it is informative to consider the working principles of current MEMS chemical sensors. Chemical microsensors often work on the basis that a change in the chemical environment induces a change in the electrical properties of the sensing material. For example, the resistance of metal oxides is sensitive to particular gases, and the change in resistance can be detected using a Wheatstone bridge. The change in resistance is proportional to the amount of target gas absorbed by the metal oxide, which in turn is proportional to the concentration of the gas. Similarly, absorption of chemical species can change the dielectric constant of an insulating material between two metal plates; the change in capacitance can be used to quantify the
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amount of this species. Many other changes in physical properties can also be used as the means for detecting chemical species, including changes in optical properties [85]. An electrochemical response can also be used for chemical sensing. For example, hydrogen peroxide is a product of the oxidation of glucose, and it can be detected by current induced at an electrode held at an electrochemical potential that oxidizes the hydrogen peroxide. The magnitude of the current is related to the glucose concentration [86]. This sensing technique can be further enhanced by the incorporation of glucose oxidase into the conducting polymer matrix [87]. Electrochemical potentials are also used to determine ion concentrations, as in the ion-selective field effect transistors (ISFET) that have been constructed from silicon for more than 30 years [88]. Most biological analyses rely on specific interactions, such as antigen–antibody binding or DNA hybridization. These binding events can be identified in a wide variety of ways. One of the most common is the detection of a fluorescence signal arising from a probe bound to the target. DNA sensing technologies using a complementary metal–oxide semiconductor (CMOS) system are reviewed elsewhere [89]. The ideal chemical and biological sensors would have the following features: high sensitivity and specificity=selectivity to the target species, linear response with no hysteresis, reversibility, no fouling, stability to environmental variations (temperature, humidity, etc.), and reproducibility. The great variety of chemical-induced property changes that can occur (e.g., color, resistivity, dielectric constant, volume, etc.) means that there is a correspondingly wide variety of sensing mechanisms.
14.3.2 Materials Selection Given the large number of sensing modalities, ideal material properties for chemical and biological sensors are not well defined. As a general rule, an electrical output is the most desirable because conversion to an electrical signal is required for signal processing (amplification, filtering, correction for environmental factors) and the implementation of recognition and quantification algorithms. Because no ideal sensors exist, arrays of sensors are used instead of single sensors in practice. The response of an array of materials not only reduces the incidence of false positives (an output indicating detection of the target when no target is present), which is a serious problem for chem-bio sensing of hazardous species, but it also allows quantification of individual components in a mixture of species.
14.3.3 Conjugated Polymer Sensors Conducting polymers are widely acknowledged as useful sensing materials for chemical and biological species [90] and have been used in electronic noses for vapor analysis [91,92]. Semiconductor materials, in general, are desirable as sensors, because environment-induced changes to the doping level or band structure can lead to large changes in electrical properties that can be easily detected using simple circuits. Electrical signals thus, generated are compatible with data acquisition, storage, and communication systems. Changes in the conductivity of conducting polymers can occur through chemical reactions, leading to changes in doping levels or through changes in polymer conformation (e.g., due to swelling caused by the absorption of chemical species). These changes are reversible, which is a key requirement for sensors. The speed of response, however, can be somewhat slow due to the need for diffusion of chemical species into the bulk of the polymer. For this reason, developing micro- or nanoscale sensors using conducting polymer gives major improvements in response times [93]. This section reviews several examples of miniature chemical sensors using conducting polymers and the performance improvements obtained. The simplest microsensor consists of a pair of electrodes covered by the sensing material and a circuit for detecting changes in resistance. An example of microscale humidity sensor with polyaniline as the sensing material is illustrated in Figure 14.7. Screen printing has been used to deposit both the conjugated polymer sensing material and the underlying carbon electrodes [56]. Wallace and coworkers [56] deposited polypyrrole with different dopants over carbon or gold electrical contacts for the detection of alcohol vapors. The carbon electrodes were screen printed on polyester substrates. They were 0.9 mm wide and 170 mm long and had 0.5 mm spacing. The gold electrodes were deposited by standard lithography techniques on silicon
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FIGURE 14.7 Construction method of a microhumidity sensor based on polyaniline: (a) photograph of completed sensor with connectors attached; (b) magnified image of platinum wire interconnects (wire diameter ¼ 125 mm); (c) magnified image of polyaniline-coated platinum wires. (From Spinks, G.M., McGovern, S.T., and Wallace, G.G., Sensor Actuator B-Chem., 107, 657, 2005.)
and were much smaller: 10 mm wide and 2 mm long with 10 mm spacing [95]. Both sensors gave similar detection limits, response times, and analyte discrimination, but results in the formation of screenprinted electrodes that were less reproducible; the authors believed, however, that results could be improved with further work. Screen printing is a low-cost fabrication method, and is attractive because the cost of sensors is one determinant of widespread adoption, particularly for disposable sensors [56]. The reduction in size of conjugated polymer chemical sensors has been shown to lead significant performance improvements. In one example, Seeber and coworkers [23] compared the electrochemical detection of ascorbic acid with both 3 mm and 10 mm diameter platinum electrodes coated with polythiophene. The microelectrode was found to give lower detection limits and to operate in poorly conductive electrolytes, which are commonly encountered in practical applications. Shrinking sensor dimensions improve response time by reducing diffusion distances. Nanofibers are, therefore, of considerable interest. Polyaniline nanofiber chemical sensors for toxic NH3 vapor have been developed using a modified electrospinning process. Typically these nanofibers are produced from a blend of a soluble conjugated polymer in a second polymer host [96–99]. Recently, modifications to produce aligned nanofibers have been reported [100–102]. Craighead and coworkers [93] deposited a single nanofiber across four gold microelectrodes (Figure 14.8a) that responded to NH3 vapor within 75 s, although it took several minutes to recover when the source of NH3 was removed (Figure 14.8b). The response time is correlated with the diffusion time for the gas to enter the fiber, and so the authors argue that nanofibers can provide faster response than micro- or macroscale systems.
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100 nm
10 µm
(a)
0.6
0.5
R/R0−1
0.4
0.3 59 ppm 0.2
42 ppm 0.8 ppm 23 ppm 8 ppm
0.1
0 (b)
0
200
400
600
800 Time (s)
1000
1200
1400
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FIGURE 14.8 (a) SEM micrograph of a single polyaniline blend nanofiber lying across four gold microelectrodes; (b) typical nanofiber response to various concentrations of ammonia vapor where the sensor was exposed to the ammonia vapor for 10 s followed by air for 2–4 min. (From Liu, H.Q., Kameoka, J., Czaplewski, D.A., and Craighead, H.G., Nano Lett., 4, 671, 2004. With permission.)
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Actuators
There are considerable potential advantages in the use of conjugated polymer actuators for MEMS [103– 105]. As described below, various prototype MEMS actuators using conjugated polymers have already been constructed with promising performances.
14.4.1 Current MEMS Actuator Principles The most common macroscopic actuators are motors. Electromagnetic motors do not operate efficiently at the microscale, but alternative actuation mechanisms become favorable. In conventional MEMS, other types of actuators have been used to generate movement, such as electrostatic actuators, thermal actuators, piezoelectric crystals, SMAs, and magnetic actuators. The most commonly used actuation mechanism in MEMS is electrostatic force generated between oppositely charged electrodes. Electrodes are fabricated so that they are tethered but can move a small distance under the attractive force. Comb drives multiply the force through numerous interdigitated fingers on the electrode. These actuators can operate at high frequency (10 kHz). To produce large displacements (up to several millimeters), a separate part is pushed by the actuators. Each displacement step of this part is small, but a large number of steps can be produced at high frequency and summed. Thermal actuators are also quite common and are of two main types: bimorphs and unimorphs. In both types the heat is generated through resistance, or Joule, heating. Bimorphs are essentially bimetallic strips, undergo out-of-plane bending on heating due to a mismatch in thermal expansion coefficients. Unimorphs are single materials with a U- or V-shape that deform in-plane on heating. The U-shape has one thin ‘‘hot’’ arm and one thick ‘‘cold’’ arm, with the hot arm expanding more so that the actuator bends. The V-shape elongates on heating, and this configuration is also used to drive moving shafts. Thermal actuators undergo large displacements but are limited in speed by their rate of cooling, although they can still achieve kilohertz frequencies. SMAs also operate through a thermal cycle, giving a volume change that can be quite large (up to 8%) [106]. Actuation occurs through a change in their crystal structure. When laminated against a second layer, the SMA can also induce a bending motion. This actuation mechanism is, however, rarely used possibly due to the difficulty of depositing and patterning these materials. Piezoelectric crystals change shape when subjected to an applied potential. The strain in ceramic piezoelectrics is typically small (0.1%), but the operating frequencies can be high (kHz). As mentioned above, quartz is one of the key members of this class of materials. Magnetic actuators have also been developed on the microscale. These are primarily devices that respond to a magnetic field created externally (off the chip), but devices have also been made, which generate the magnetic field on-chip by running a current through a coil. There are fewer of the latter type because of the difficulty of fabricating coils on a two-dimensional surface.
14.4.2 Materials Selection For actuators, the key performance metrics are stress and strain. For autonomous applications, low power is also of importance, and for some devices, fast response times and long cycle lives are required. The most important mechanical property is the elastic (Young’s) modulus, because it determines the strain and stress that can be produced. Large strains can be achieved by materials with low modulus, but at the cost of force, and vice versa for materials with high modulus. In materials with high modulus, large strokes are only possible with bending configurations. Thermal expansion, creep, yielding, and fracture all interfere with the mechanical response.
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The actuator materials currently used for MEMS have been found to be suitable in all cases except one: actuators for locomotion, gripping, and other interactions with the environment. Actuators made from inorganic materials face two key stumbling blocks in this arena: large footprints and brittleness. As pointed out in 1992 by Elwenspoek et al. [107], microrobots are still science fiction because there are no actuators useful for locomotion (and also no suitable miniature power supplies), and this situation has not changed in intervening years. Conventional actuators also suffer from relatively large ‘‘footprint’’ and this large size limits the degree of miniaturization and actuator densities that can be achieved. In contrast, arrays of polymer thermal actuators [108] have been fabricated that have a high actuator density and can be used to exploit parallel arrangements, such as for the handling of large objects. The brittleness of inorganic materials means that they readily break upon contact with macroscale objects. The polysilicon legs used in the thermal actuator array above snapped when too much weight was applied. Similarly, polysilicon microgrippers have been known to break if touched by a macro-object or if exposed to air velocities higher than 1 m=s [109]. In contrast, polymer actuators are compliant. Figure 14.9 shows a chip with polymeric thermally actuated bending legs [110]. This device could withstand when dropped from a height of 50 cm, as well as when pushed down onto the surface without damage and having the legs return to their original position. Figure 14.10 shows a conjugated polymer bending microactuator that was hit by a macroscale object and then operating normally afterward [35]. These advantages have given polymer actuators a significant edge, putting them at the forefront for microrobotics applications. The device in Figure 14.9 by Ebefors was the first, and is still the only walking microdevice. As will be discussed later, the only microrobot arm ever built had conjugated polymer actuators. There are a large number of polymer actuation mechanisms, each suitable for particular applications. The most ubiquitous for MEMS is thermal actuation, which is accomplished with conventional polymers and which makes use of differences between the CTE of the polymer and a second layer, which is often a metal. The actuator is heated by running current through the metal layer. To give some example metrics [108,111], stress can be up to 80 MPa and lifetimes greater than 108 cycles. Strains are up to 2%, but decrease above 1 Hz. They have intrinsically high power consumption because they turn the input energy into irrecoverable heat.
FIGURE 14.9 Chip with V-groove polyimide thermally actuated legs (left). Cross-section schematic of one of the actuators (right). (From Ebefors, T., Polyimide V-groove joints for three-dimensional silicon transducers. Department of Signals, Sensors, and Systems. Royal Institute of Technology (KTH), Stockholm, Sweden, 2000.)
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FIGURE 14.10 Conjugated polymer microactuator being hit by a macro-object (top row). Actuator working normally afterward (bottom row). (From Smela, E., J. Micromech. Microeng., 9, 1, 1999. With permission.)
As mentioned above, thermal actuator based on a phase change in inorganic SMA materials has also been microfabricated. Polymer analogs of these include polydiacetylene [112], in which the side chains change configuration, and liquid crystal elastomers [113], which undergo a nematic-to-isotropic transition. They have strains of up to 400%, but stresses of only 0.1 MPa as the Young’s modulus is only about 105 GPa. Response times are about 1 Hz, with actuation by Joule heating of embedded conducting particles. Efficiencies in converting electrical energy to mechanical energy are low, as for other thermal actuators. Because there is a phase transition involved, the actuator switches between two endpoint positions: it is not continuously adjustable. No polymer SMA actuators have yet been microfabricated. Electronic electroactive polymers (EAP) include dielectric elastomers [114], piezoelectric polymers, such as polyvinylidenefluoride (PVDF) [115], and electrostrictives. Actuation in these materials is based on electric fields, so high voltages are required. These actuators have, in general, high strength and speed, and the actuation mechanisms are fairly well understood. For PVDF, the strains are small, approximately 0.1%, but stresses are 100 MPa and the Young’s modulus is 2 GPa. They can hold any intermediate position as the strain is linear with the electric field. Response times can be greater than 105 Hz, lifetimes are high, and efficiencies are an impressive 20%. However, because of the required poling step, these materials are difficult to process and have not been widely adopted for MEMS. Dielectric elastomers have strains of up to 300%, stresses of 10 MPa, a modulus of 104 to 102 GPa, and response times of kilohertz. They also hold intermediate positions, but the strain is proportional to the square of the electric field. Lifetimes in excess of 107 cycles have been demonstrated. These have not yet been microfabricated, but it should be possible. Conjugated polymers belong to the family of ionic EAPs, in which actuation is based on ion and solvent transport. As a consequence, all require the presence of an electrolyte and are relatively slow, although at the microscale their speed improves as ion transport lengths decrease. They can be actuated at low voltages, which makes it possible to drive them using standard on-chip circuitry, but they consume high currents, so efficiencies are poor. Other members of the family of ionic EAPs include ionic polymer–metal composites (IPMCs) [116], gels (which are not actually electroactive), and carbon
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nanotube papers (which are not actually polymeric). Of these, only gel actuators have been microfabricated to date, including microstructured thermally [117] and optically [118] activated gels. For comparison, the key metrics of these materials are summarized here. IPMCs are estimated to have a strain of a few percent, but only bending movements are possible. Their moduli are approximately 0.3 GPa, the response times of mesoscale actuators are 1 s, and the actuators are robust, but lifetimes have not been reported. In most electrolytes, these actuators cannot hold a fixed position. A gel microactuator is described in Section 14.5.1. Strains in gels are typically tens of percent up to 1000%, but the Young’s modulus is low, about 0.1–1 MPa. Response times are diffusion limited, but moderate speeds (10 s) are possible on the microscale. As they are not electrically controlled, there is no efficiency metric. Carbon nanotube sheets [119] are electrochemically actuated like conjugated polymers. Strains are a few percent, and the modulus is 1 GPa. Response times are approximately 1 Hz.
14.4.3 Conjugated Polymer Microactuators Conjugated polymer microactuator devices have been reviewed many times previously [120–122], and so this section will give only a brief. For further information, refer to those reviews and the publications referenced therein. Conjugated polymer microactuators have primarily been bending bilayers. These were first demonstrated in 1993 [123], and a robust process sequence that incorporated rigid plates was shown in 1995 [124]. In order to avoid exposing the polymer to potentially damaging etchants, allow the formation of larger structures without etch holes and eliminate the large step between the free and anchored regions associated with using a sacrificial layer for the release of the structures, a ‘‘differential adhesion’’ technique was developed to free the bilayers [124]. The process uses a chromium metal layer as an adhesion layer between the structural gold layer and an oxidized silicon substrate (or more preferably a Cr–Au layer). By patterning the adhesion layer before deposition of the gold, the adhesion level could be controlled across the substrate surface because gold does not adhere to Si or SiO2. Subsequent deposition of polypyrrole and other structural polymers completed the device. When the PPy was electrochemically cycled the volume change produced sufficient bending stresses to allow the flap to break free from the poorly adhered areas. The hinge, however, remained attached to the substrate at the regions where chromium was deposited to provide strong adhesion. In 2000, a microrobot arm was demonstrated that had multiple independently controlled actuators [125]. All of these actuators were operated when immersed in a liquid electrolyte (Figure 14.11). The metrics for conjugated polymer actuators show a good balance between stress and strain. Strain is highly anisotropic, with reported values of 3% in-plane and 40% out-of-plane. Strain is also strongly thickness-dependent in the thin films used in microstructures, and the existence of strain gradients must be taken into account when designing devices [126]. Isotonic stresses in polyaniline fibers have been reported to be as high as 34 MPa, though strain decreases with stress [127–129]. Blocked stresses are considerably lower. Modulus measurements for polypyrrole are 0.2–4.0 GPa [130], depending on oxidation level [131] and deposition conditions. Response times for microactuators are 1 Hz, with this metric strongly dependent on thickness. The actuators readily hold a fixed position, and hysteresis can be virtually eliminated by controlling charge, rather than voltage. Polypyrrole degrades in aqueous electrolytes, with no electroactivity present in 1 mm thick PPy films doped with dodecylbenzenesulfonate PPy(DBS) after 40,000 cycles. However, in ionic liquids a million cycles have already been demonstrated in polyaniline [132]. (This value does not present an upper limit; just when testing was halted.) Because polypyrrole operates in aqueous electrolytes at room temperature, the largest niche for conjugated polymer microactuators is biomedical applications. Commercialization efforts are underway for blood vessel connectors, a valve to prevent urinary incontinence, and a Braille display [25,122,133]. One area that requires further research is the temperature-dependence of actuator metrics, because for biomedical applications the devices must be operated at 378C. In PPy(DBS) microactuators, strain increases from room temperature to body temperature by 45%, and they are 250% faster, but the blocked force drops [126].
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Finger actuator
Wrist actuator
PPy
Gold Elbow actuator (a)
(b)
(c)
FIGURE 14.11 Microrobot arm with PPy and gold on SU8. (a) The arm has three sets of actuators. (b) The elbow and wrist actuators are bent to place the fingers over a small bead. (c) The finger actuators grab the bead and the wrist actuator unbends so that the arm holds the bead face-up. (From Jager, E.W.H., Ingana¨s, O., and Lundstro¨m, I., Science, 288, 2335, 2000.)
The next step for this technology is the development of ‘‘dry’’ microactuators that can work outside of a liquid. On the macroscale, it has been demonstrated that conjugated polymers can actuate in ionic liquids [132,134]. These room temperature liquid salts can be incorporated into gels [135], which will allow the fabrication of patterned electrolyte containing solid layers using photolithographic techniques. This development will expand the range of possible applications. The low efficiency of these actuators will limit the applications, however. Less than 1% of the input electrical energy is converted to mechanical work [35]. The commercialization efforts have focused on devices for which efficiency is not an issue. The urinary incontinence microvalves are small enough, and are actuated infrequently enough, to run off of a watch battery for over a month, the blood vessel connectors require only actuation at the time of insertion, and the display would typically be plugged in, like a laptop, except for short periods of time. An autonomous microrobot, however, would consume too much power for these actuators. It is important to keep in mind that an electron is consumed for each ion transported, so that a pencil-sized actuator switching in 1 s would draw a kiloamp of current [136].
14.5
Microfluidic Systems
Microfluidic systems are developed to encapsulate an entire chemical or biological assay into one compact unit. This requires sample uptake, sample preconditioning, and analysis capability to be built into the system.
14.5.1 Current Microfluidics: Principles and Methods To move fluids through the system, electrical fields are used in almost all cases, exploiting electroosmosis, electrophoresis, and other voltage-driven mechanisms that can be used at the microscale, but not at the macroscale. These methods have the advantage of having no moving parts. However, in some
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cases electric fields cannot be used, and mechanical pumps with moving parts are required. Also, valves are most often mechanical. Microvalves have been demonstrated virtually with all of the actuation mechanisms described above. Further details are available elsewhere [5,11].
14.5.2 Materials Selection Ionic electroactive polymers are well suited to microscale applications because the ion-transport limitations that result in unacceptably low speeds on the macroscale are lifted. A good example of the favorable scaling is given by microfabricated gel actuators. On the macroscale, these actuators can take hours, because the volume change is diffusion-limited. However, diffusion times go as length squared, so improvements in response times upon microfabrication are dramatic. Beebe et al. [137] fabricated gelcoated posts within microchannels. The 100 mm thick gel skins expanded upon exposure to a step change in pH to pinch off a microchannel. It took 8 s for the expansion to reach half-maximum. Conjugated polymers, in which ion transport is dominated by migration (i.e., movement under an electric field) under a step potential [138], respond much more quickly, and are thus, even more attractive.
14.5.3 Conjugated Polymer Pumps and Valves There have been several applications of conducting polymers to microfluidics. The fabrication protocols for microfluidic devices are amenable to the use of conducting polymers, as outlined above, and conjugated polymers could potentially be used both as sensors and as actuating elements (pumps, valves) in lab-on-a-chip devices for chemical analysis. Given the potential advantages, it is likely that the number of microfluidic devices using conjugated polymers will increase soon. The prototype conducting polymer devices that have been reported are surveyed in this section. Pettersson et al. [139] fabricated a valve based on polypyrrole bending bilayers attached to rigid polymer plates. They fabricated the actuators on a flat surface, and then bonded that to a polymer cover with the fluid channels. Fluid flow was reduced when the plates were rotated perpendicular to the surface so that they blocked the channel. A similar device has been developed to operate a flap that opens and closes the cover on a ‘‘microvial’’ for cell clinic processes [140–142]. Similar polypyrrole bilayers have recently been used by Park et al. [143] to separate embryonic cells for automated lab-on-a-chip cell manipulation. Embryo cells were placed in a reservoir and moved into an injection port with a PPy(DBS) valve. As the first cell passed the valve, it was closed to prevent other cells from entering the following manipulation chamber. This device integrated four actuation technologies: pressure-driven flow, dielectrophoretic cell rotation, precise positioning through suction, and PPy valves. The fabrication process for the valves was the one previously developed by Smela [35] utilizing differential adhesion. Subsequent to gold deposition, microchannels were fabricated from SU8, and then PPy was deposited in a photoresist template. The channel was sealed by bonding to a glass cover using a UV-curable adhesive. The PPy valves were normally closed without an applied potential, and they were opened upon reduction. There were several benefits cited for using the PPy valves: biocompatibility, small size, amenability to fabrication within a microchannel, and good cell viability through avoidance of lasers and electric fields. Lee et al. [144] presented a microfabrication process for a membrane pump using polypyrrole as the actuating material. The diaphragm displacement in liquid electrolyte was hundreds of micrometers. Although further work to improve reliability is required, this is a promising development. More recently, Berdichevsky and Lo [145] microfabricated a microvalve that made use of the large out-of-plane strain in PPy(DBS). Microchannels were fabricated from PDMS, which is an elastomer, by soft lithography. Polypyrrole was deposited onto a gold-working electrode adjacent to a gold counter electrode in a large channel. The gold-working electrode had the shape of a post to prevent delamination of the PPy. A PDMS membrane topped this channel, and a second, narrower, hemispherical microchannel was fabricated directly over the PPy post. Upon expansion of the PPy, the membrane pushed upward into the second channel, closing it off to form a complete seal.
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Polpyrrole Stainless steel auxiliary electrode mesh Porous separator containing electrolyte
Outlet
(a) Flexible polyurethane Polypyrrole tubing working electrode 0s CP Pump 30 s (b)
FIGURE 14.12 (a) Schematic illustration of the design of a conducting polymer-based micropump; (b) fluid transport achieved with the first prototype device. (Courtesy of Y. Wu, University of Wollongong.)
Wu and coworkers [146] have fabricated a prototype mesoscale pump. The device consisted of concentric cylinders of conducting polymers with a porous separator containing the electrolyte (Figure 14.12a). When a voltage was applied, the volume change of the conducting polymer inner cylinder compressed a flexible tube containing the fluid. The resultant pressure caused the fluid to flow (Figure 14.12b). Using standard microfabrication techniques, it should be possible to produce smaller pumps integrated directly into a microchannel. Changes in polymer surface energy have been described as potentially useful in microfluidic applications [147,148]. For example, by controlling the surface energy it is possible to move fluids across the surface. Conducting polymers have been investigated as active surfaces for switchable surface energy [149,150]. In one study [150], water contact angles on both polyaniline (DBSA doped) and poly(3-hexyl thiophene) were investigated in the oxidized and the reduced states. The surface energy was found to change significantly, as evidenced by changes in contact angles of up to 408. This effect was then exploited in a cell constructed so as to produce a gradient of doping level in a polyaniline film. This gradient induced the movement of a water droplet toward the reduced side. In a separate study [149], polypyrrole doped with DBS was coated onto microchannels formed in PMMA. By applying a reducing potential to the polypyrrole, electrolyte was made to move along the channels. A microfluidic sensing application has been reported by Ghodssi and coworkers [151], who electrodeposited chitosan, a biopolymer, on polypyrrole tracks. Polypyrrole was specifically chosen because of its biocompatibility. The chitosan was functionalized to bind with various biomolecules.
14.6
Energy Storage Systems
Whereas microsystems can be run off of a small watch battery, for further miniaturization an on-chip energy source is required. Microbatteries and microcapacitors are also envisaged to be useful power sources for memory backup, active cables (battery is embedded in the cable), and radio frequency identification devices [152,153].
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14.6.1 Current Microfabricated Energy Storage Methods ‘‘Power MEMS’’ [6] is the term used to describe miniaturized sources of electrical power. Batteries remain as the major source of portable electrical power, but MEMS-type turbines [6,154] and fuel cells are investigated as alternatives. The main limitations of batteries, for autonomous MEMS applications such as microrobots, are their low energy and power densities. Turbines and fuel cells have higher densities, but these systems are not yet available on the microscale. Fuel cells have, however, been miniaturized to some extent, with the size of a pack of cards already available commercially [154] for applications like powering cell phones.
14.6.2 Materials Selection Because of the diverse nature of producing electrical energy from these different systems, it is not possible to clearly specify the necessary materials properties. However, the relatively simple microfabrication techniques available for conjugated polymers may enable the development of microscale power sources, with efficiencies proving to be competitive. On a macroscopic scale, conjugated polymers have also been extensively researched for batteries, capacitors, and photovoltaics [155].
14.6.3 Conjugated Polymer Batteries and Capacitors Both batteries and electrochemical capacitors consist of two electrodes on either side of an electrolyte. Whereas batteries are traditionally considered to generate current through Faradaic redox reactions occurring at the electrodes, labeled anode and cathode, electrochemical capacitors store charge solely in a double layer at the two electrode plates. In the former, electrons are transferred to a species in the electrolyte to the electrode, whereas in the latter charge is merely redistributed within each phase. In CPs, the Faradaic reaction consists of pulling an electron off the backbone, or adding one, and in response an ion moves to or away from the chain. Capacitive current occurs in response to applying a potential to the polymer and charging the double layer, without changing the net charge on the backbone. However, at the molecular scale it is impossible to clearly distinguish between Faradaic and capacitive reactions in a conjugated polymer. In an oxidized state, the polymer is a porous electrode with a large surface area, and thus has a significant capacitance. Thus, at the nanoscale the distinction between batteries and electrochemical capacitors becomes blurred, and only the charge and discharge rates are distinguished between ‘‘battery’’ and ‘‘capacitor’’ behavior. A key performance criterion for both systems is to store as much electrical energy per unit mass as possible, and another is how fast that energy can be delivered. Sung et al. [153,156] have described the fabrication and performance of prototype microcapacitors built with PPy and polythiophene electrodes. Interdigitated gold electrodes were patterned on silicon using standard photolithography, and the conducting polymers were then electropolymerized across them. The capacitor performance was evaluated with different polymer thicknesses and electrolytes. A subsequent development incorporated solid polymer electrolytes to produce a solid-state device [153], although the performance was diminished compared with devices operated in liquid electrolytes. Microbatteries are typically produced from multiple thin layers. An alternative design has been reported in which polypyrrole was proposed to be used as the cathode and lithiated carbon as the anode [152]. The performance targets specified were: a cell voltage of 3–4 V; a discharge rate of 0.1 mA=cm2, and a 200 cycle lifetime (at 80% discharge). Although the conceptual design was published in 1999 and the properties of the carbon anode considered, no further work has been reported on this system and no CP microbatteries have yet been constructed. The advantage of using conducting polymers for power generation lies in potentially low-cost manufacture. However, it remains to be demonstrated whether conjugated polymer batteries and
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capacitors that will produce practically useful energy and power densities can be developed. The lack of activity in this area in recent years suggests that the challenges are considerable.
14.7
Future Challenges and Possibilities
Earlier sections have demonstrated the wide applicability and advantages of conjugated polymers in microsystems. However, continued research is needed to fully assess the merits of conducting polymers in MEMS, to quantify their performance capability, and to understand the underlying sensing and actuation mechanisms. For some of these devices, substantial improvements are still necessary to give them the edge over competing technologies. Nevertheless, the use of conjugated polymers in prototype devices has increased rapidly in the last 5 years as the potential benefits have become clearer. In addition, there are reasons to believe that the performance of conjugated polymer systems will improve as dimensions shrink, decreasing ion transport times and lowering resistive losses. It is difficult to predict which of the prototype devices developed in research laboratories will break through to major commercial production. As for any other technology, whether the prototype devices move onto commercialized products depends on their performance, and cost versus competing technologies. A review of the history of the development of silicon MEMS shows that production of single devices for mass markets has been the dominant pattern in commercialization. The first products were pressure sensors for the automobile mass market. Next came the accelerometers for the same market, and more recently optical switches and ink-jet printheads. Similar parallels can be seen in the development of conducting polymer devices. Of the wide variety of prototype sensors, actuators, and electrical components that have been demonstrated over the past 10–15 years, only polymer LEDs have been completely commercialized. Smart systems utilizing MEMS and conjugated polymers, which were anticipated to be the ‘‘killer app,’’ have yet to make their appearance. The future application of conjugated polymers in microsystems will depend largely on the unpredictable needs of the market, but will be influenced by improvements in device performance (including speed, reliability, efficiency, and lifetime) and cost (determined by both base material costs and fabrication costs). Decisions to invest in this new technology will be based on a balance of potential pay-offs against the risks and cost of the investments. Realities like the high power requirements for actuators will be ‘‘show-stoppers’’ for some applications. Fortunately, there do not appear to be insurmountable problems in adapting conjugated polymer to microsystems, as evidenced by the number of successful prototypes already constructed. These prototypes have also begun to tackle the packaging issues that can plague commercialization: how to make reliable electrode connections, prevent delamination, encapsulate the electrolyte, and integrate reference electrodes. The continued miniaturization of electronic devices has already moved them into the realm of nanotechnology, with dimensions less than 100 nm. Likewise, MEMS have also begun the move into nanoelectromechanical systems (NEMS). As the ‘‘top-down’’ approaches of traditional microfabrication techniques approach a lower limit, alternative ‘‘bottom-up’’ approaches are increasingly utilized to assemble devices at the molecular level. Nanofabrication techniques have recently been reviewed by Whitesides and coworkers [157]. Sensors have already been produced using individual conjugated polymer nanoparticles [93]. The development of such working prototypes at the nanoscale will require that fabrication and connection issues be addressed. In addition, valuable information can be obtained regarding the performance of conjugated polymers when operated at the nanoscale. Quite a new phenomenon may emerge, such as the discrete electrochemical switching described above [26]. However, it is likely to take some considerable time before nanofabrication techniques can match the mass production capacity and low cost of microtechnology. In the short to medium term, NEMS is likely to remain as a fruitful and fascinating research activity. Looking back on the history of MEMS commercialization, we can confidently predict that the commercial impact of NEMS will be highly unpredictable!
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115. Zhang, Q., and J. Scheinbeim. 2001. Electric EAP. In Electroactive polymer (EAP) actuators as artificial muscles: Reality, potential, and challenges, ed. Y. Bar-Cohen, 89–138. Bellingham: SPIE Press. 116. Shahinpoor, M., and K.J. Kim. 2001. Ionic polymer–metal composites. I. Fundamentals. Smart Mater Struct 10:819–833. 117. Wang, J., Z. Chen, M. Mauk, K.-S. Hong, M. Li, S. Yang, and H.H. Bau. 2005. Self-actuated, thermo-responsive hydrogel valves for lab on a chip. Biomed Microdevices 7:313–322. 118. Li, M.-H., P. Keller, B. Li, X. Wang, and M. Brunet. 2003. Light-driven side-on nematic elastomer actuators. Adv Mater 15:569–572. 119. Baughman, R.H., C. Cui, A.A. Zakhidov, Z. Iqbal, J.N. Barisci, G.M. Spinks, G.G. Wallace, A. Mazzoldi, D. de Rossi, A.G. Rinzler, O. Jaschinski, S. Roth, and M. Kertesz. 1999. Carbon nanotube actuators. Science 284:1340–1344. 120. Baughman, R.H. 1996. Conducting polymer artificial muscles. Synthetic Met 78:339–353. 121. Jager, E.W.H., E. Smela, and O. Inganas. 2000. Microfabricating conjugated polymer actuators. Science 290:1540–1545. 122. Smela, E. 2003. Conjugated polymer actuators for biomedical applications. Adv Mater 15:481–494. 123. Smela, E., O. Inganas, Q.B. Pei, and I. Lundstrom. 1993. Electrochemical muscles—micromachining fingers and corkscrews. Adv Mater 5:630–632. 124. Smela, E., O. Inganas, and I. Lundstrom. 1995. Controlled folding of micrometer-size structures. Science 268:1735–1738. 125. Jager, E.W.H., O. Ingana¨s, and I. Lundstro¨m. 2000. Microrobots for micrometer-size objects in aqueous media: Potential tools for single cell manipulation. Science 288:2335–2338. 126. Christophersen, M., and E. Smela. (in press). Characterization and modeling of PPy bilayer microactuators. Part 4: Effect of temperature. Sensor Actuator B. 127. Smela, E., W. Lu, and B.R. Mattes. 2005. Polyaniline actuators, Part 1: PANI(AMPS) in HCl. Synthetic Met 151:25–42. 128. Spinks, G.M., L. Liu, D. Zhou, and G.G. Wallace. 2002. Strain response from polypyrrole actuators under load. Adv Funct Mater 12:437–440. 129. Spinks, G.M., and V.-T. Truong. 2005. Work-per-cycle analysis of electromechanical actuators. Sensor Actuator A-Phys 119:455–461. 130. Bay, L., S. Skaarup, K. West, T. Mazur, O. Joergensen, and H.D. Rasmussen. 2001. Properties of polypyrrole doped with alkylbenzene sulfonates. Presented at the Proceedings of the SPIEs 8th International Symposium on Smart Structural Materials, Electroactive Polymer Actuators and Devices (EAPAD), Newport Beach, CA. 131. Murray, P., G.M. Spinks, G.G. Wallace, and R.P. Burford. 1998. Electrochemical induced ductilebrittle transition in tosylate-doped (pTS) polypyrrole. Synthetic Met 97:117–121. 132. Lu, W., A.G. Fadeev, B.H. Qi, E. Smela, B.R. Mattes, J. Ding, G.M. Spinks, J. Mazurkiewicz, D.Z. Zhou, G.G. Wallace, D.R. MacFarlane, S.A. Forsyth, and M. Forsyth. 2002. Use of ionic liquids for pi-conjugated polymer electrochemical devices. Science 297:983–987. 133. Immerstrand, C., K. Holmgren-Peterson, K.E. Magnusson, E. Jager, M. Krogh, M. Skoglund, A. Selbing, and O. Inganas. 2002. Conjugated-polymer micro- and milliactuators for biological applications. MRS Bull 27:461–464. 134. Vidal, F., C. Plesse, and D. Teyssie. 2004. Long-life air working conducting semi-IPN=ionic liquid based actuator. Synthetic Met 142:287–291. 135. Fuller, J., A.C. Breda, and R.T. Carlin. 1998. Ionic liquid-polymer gel electrolytes from hydrophilic and hydrophobic ionic liquids. J Electroanal Chem 459:29–34. 136. Madden, J.D.W., P.G.A. Madden, and I.W. Hunter. 2002. Conducting polymer actuators as engineering materials. Presented at the Proceedings of the SPIEs 9th International Symposium on Smart Structural Materials, Electroactive Polymer Actuators and Devices (EAPAD), San Diego, CA.
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137. Beebe, D.J., J.S. Moore, J.M. Bauer, Q. Yu, R.H. Liu, C. Devadoss, and B.-H. Jo. 2000. Functional hydrogel structures for autonomous flow control inside microfluidic channels. Nature 404:588– 590. 138. Wang, X., B. Shapiro, and E. Smela. 2004. Visualizing ion transport in conjugated polymers. Adv Mater 16:1605–1609. 139. Pettersson, P.F., E.W.H. Jager, and O. Ingana¨s. 2000. Surface micromachined polymer actuators as valves in PDMS microfluidic system. Presented at IEEE-EMBS Special Topic Conference on Microtechnologies in Medicine and Biology, Lyon, France. 140. Jager, E.W.H., C. Immerstrand, K.H. Petersson, K.-E. Magnusson, I. Lundstro¨m, and O. Ingana¨s. 2002. The cell clinic: Closable microvials for single cell studies. Biomed Microdevices 4:177–187. 141. Liu, Y., N.M. Nelson, P. Abshire, and E. Smela. 2004. Biolab-on-a-chip for capturing, culturing, and in-situ investigation of living cells. Presented at MicroTAS 2004, Malmo¨, Sweden. 142. Jager, E.W.H., E. Smela, O. Ingana¨s, and I. Lundstro¨m. 1999. Applications of polypyrrole microactuators. Presented at SPIEs 6th International Symposium on Smart Structural Materials, Electroactive Polymer Actuators and Devices (EAPAD), Newport Beach, CA. 143. Park, J., S.-H. Jung, Y.-H. Kim, S.-K. Lee, and J.-O. Park. 2005. Design and fabrication of an integrated cell processor for single embryo cell manipulation. Lab Chip 5:91–96. 144. Lee, S.-K., S.-J. Lee, H.-J. An, S.-E. Cha, J.K. Chang, B. Kim, and J.J. Pak. 2002. Biomedical applications of electroactive polymers and shape memory alloys. Presented at Smart Structures and Materials, Electroactive Polymer Actuators and Devices, San Diego, CA. 145. Berdichevsky, Y., and Y.-H. Lo. 2003. Polymer microvalve based on anisotropic expansion of polypyrrole. Presented at the Material Research Society Symposium. Fall 2003 Meeting, Boston. 146. Wu, Y.Z., G.G. Wallace, and G.M. Spinks 2005. The TITAN polypyrrole micropump. Smart Materials and Structures. 14:1511–1516. 147. Someya, T., A. Dodabalapur, A. Gelperin, H.E. Katz, and Z. Bao. 2002. Integration and response of organic electronics with aqueous microfluidics. Langmuir 18:5299–5302. 148. Zhao, B., J.S. Moore, and D.J. Beebe. 2002. Principles of surface-directed liquid flow in microfluidic channels. Anal Chem 74:4259–4268. 149. Causley, J., S. Stitzel, S. Brady, D. Diamond, and G. Wallace. 2005. Electrochemically-induced fluid movement using polypyrrole. Synthetic Met 151:60–64. 150. Isaksson, J., C. Tengstedt, M. Fahlman, N. Robinson, and M. Berggren. 2004. A solid-state organic electronic wettability switch. Adv Mater 16:316–320. 151. Kastantin, M.J., S. Li, A.P. Gadre, L.-Q. Wu, W.E. Bentley, G.F. Payne, G.W. Rubloff, and R. Ghodssi. 2003. Integrated fabrication of polymeric devices for biological applications. Sensor Mater 15:295–311. 152. Kinoshita, K., X. Song, J. Kim, and M. Inaba. 1999. Development of a carbon-based lithium microbattery. J Power Sources 81–82:170–175. 153. Sung, J.H., S. Kim, and K.H. Lee. 2004. Fabrication of all-solid-state electrochemical microcapacitors. J Power Sources 133:312–319. 154. Freedman, D.H. 2004. Power on a chip. Technol Rev November: 49–52. 155. Brabec, C.J., N.S. Sariciftci, and J.C. Hummelen. 2001. Plastic solar cells. Adv Funct Mater 11:15–26. 156. Sung, J.H., S.J. Kim, and K.H. Lee. 2003. Fabrication of microcapacitors using conducting polymer microelectrodes. J Power Sources 124:343–350. 157. Gates, B.D., Q. Xu, M. Stewart, D. Ryan, C.G. Willson, and G.M. Whitesides. 2005. New approaches to nanofabrication: Molding, printing, and other techniques. Chem Rev 105:1171– 1196.
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15 Corrosion Protection Using Conducting Polymers 15.1 15.2 15.3
Introduction..................................................................... 15-1 Aqueous Corrosion and Passivation.............................. 15-4 Traditional Methods of Corrosion Control .................. 15-6 Coatings . Corrosion Inhibitors Protection
15.4
.
Cathodic and Anodic
Techniques for Studying Corrosion............................. 15-10 Background . Global Electrochemical Methods . Local Electrochemical Methods . Surface Spectroscopy and Imaging Methods
15.5
Conducting Polymers for Corrosion Control—General Considerations ............................... 15-18 Electronic Interactions . Chemical Interactions . Redox Reactions at the Conducting Polymer Surface . Conducting Polymer as an Oxidant . Neutral (Undoped) Forms of Conjugated Polymers . Role of the Dopant Ion . Influence of CPs on Overall Coating Impedance
15.6
Approaches to Forming Conjugated Polymer Coatings .......................................................... 15-28 Functionalized Solvent-Processable Polymers . Conducting Polymer Composites and Blends . Electrodeposition of Conducting Polymers
15.7
Oxide Layers and Active Coatings ............................... 15-35 Physical and Electronic Properties of Oxides of Metal Surface Preparation
15.8
Dennis E. Tallman and Gordon P. Bierwagen
15.1
Methods
Conducting Polymers for Corrosion Protection—Recent Results .......................................... 15-38 Corrosion Protection of Iron and Steel Protection of Aluminum and Its Alloys Protection of Other Metals
15.9
.
. .
Corrosion Corrosion
Summary and Prognosis............................................... 15-42
Introduction
The corrosion of a material may be defined as the irreversible reaction of the material with its environment, usually resulting in degradation of the material and its properties. In this chapter, we will be concerned only with the corrosion of metals in aqueous environments under near-ambient 15-1
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conditions, and we will ignore topics such as high-temperature corrosion, erosion corrosion, environmentally induced cracking, and nonaqueous corrosion. The corrosion of a metal may be viewed as extractive metallurgy in reverse. Considerable energy must be expended to extract and purify a metal such as Fe or Al from its ore (i.e., its oxide forms) and, thus, corrosion is the thermodynamically driven process by which the metals revert to their oxide forms. For example, the energy obtainable from the oxidation of 1 mole (55.85 g) of pure Fe to g-FeOOH is 81 kcal, sufficient to power a 100 W light bulb for nearly 1 h. Thus, it is virtually impossible to completely stop such a thermodynamically favorable process. As long as we continue to use active (i.e., corrodible) metals for the construction of bridges, buildings, automobiles, airplanes, ships, industrial reactors, and other objects, there will be corrosion problems and associated costs. It is estimated that the annual cost of corrosion and its control for a developed country is approximately 3%–4% of the country’s gross domestic product; for the United States alone, that amounts to $300 billion per year [1]. As there is not much that can be done to alter the thermodynamics of corrosion of structural (active) metals in natural environments, corrosion control strategies typically focus on slowing the rate of the corrosion process. A very common corrosion control strategy is to apply one or more layers of a coating to the metal. Such a coating may simply serve as a barrier between the metal and its environment, retarding the rate at which water, oxygen, or ions from the environment reach the metal surface. On the other hand, a coating may function as more than just a barrier. The coating may be an active coating in the sense that it contains or consists of a material than can interact chemically or electrochemically with the metal, altering its corrosion behavior. For example, coatings containing hexavalent chromium are often used to control the corrosion of aluminum alloys [2], whereas coatings containing Zn particles are commonly used on steel [3]. Such coatings function through (often complex) electrochemical or chemical interactions of the active species (Cr(VI) or Zn) and their reaction products with the metal surface. By this definition, conducting polymer coatings are active coatings. There are several issues that motivate the search for new and improved corrosion control coatings. First, there is the desire to prolong the lifetime of coating systems, reducing the cost associated with maintenance, stripping, and recoating. Then, there is the need to eliminate toxic components from corrosion control coatings. For example, the Environmental Protection Agency (EPA) and the Occupational Safety and Health Agency (OSHA) will continue to restrict the use of Cr(VI) as a corrosion control component because of its toxicity and issues related to safe handling and disposal [4]. As a final motivation, there may be the opportunity to design multifunctional coatings that not only protect against corrosion, but also provide an indication of impending coating failure or corrosion onset. Such self-sensing coatings may be thought of as ‘‘intelligent (or smart) coatings’’ and represent a long-term goal of research in our laboratory. This chapter will focus on the corrosion protection of active metals (principally iron, aluminum, and their alloys) using conducting polymers. The term conducting polymer (CP) will be used to describe a highly conjugated organic polymer where some degree of electronic conductivity occurs. Other types of conducting polymers not considered in this chapter include redox polymers (where conductivity occurs by electron hopping between specific spatially and electrostatically isolated electrochemically active sites within the polymer [5]) and ionically conducting polymers (e.g., ion exchange polymers; Ref. [6], p. 208). Figure 15.1 shows examples of the more common conjugated polymers in their neutral (or undoped) state. The conductivity of these neutral polymers is typically quite low, ranging from insulating to poorly semiconducting. After partial oxidation (often referred to as p-doping), such polymers become more highly conducting, with metal-like conductivities in some cases. The partial reduction (or n-doping) of certain conjugated polymers is also possible [7,8], although most conjugated polymers are difficult to n-dope and such polymers often are not air stable. To our knowledge, there have been no reports of the use of n-doped polymers in corrosion protection, and such polymers will not be considered in this chapter. It should be noted, however, that n-doped conducting polymers could (in principle) function as sacrificial coatings in a manner similar to Zn-rich coatings [3] or perhaps as oxygen scavengers similar to sulfite and hydrazine inhibitors, and this is a research area which our group intends to explore.
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Corrosion Protection Using Conducting Polymers
NH N H
n
n
Polypyrrole
Polyaniline
S n
n
Polythiophene
FIGURE 15.1 coatings.
Poly(p-phenylene vinylene)
Neutral (undoped) forms of parent-conjugated polymers that have been studied as corrosion control
Figure 15.2 shows the interconversion between the neutral and oxidized forms of a CP, in this case polypyrrole (PPy). In this example, a mobile anion is incorporated into the oxidized polymer to maintain charge balance. This anion is expelled when the polymer undergoes reduction. If the anion is not mobile (e.g., a polymeric anion or an anionic group covalently attached to the polymer backbone), then the cations will move into the polymer upon reduction. In many cases, upon reduction, a mixture of anion expulsion and cation uptake is observed. It is also important to note that the oxidized forms of the polymers of Figure15.1 are all oxidants toward virtually any active metal [9]. That is, their equilibrium potentials are significantly positive of that of the active metals. Several possible interactions between an active metal and the CP are suggested by the preceding discussion. As noted above, the oxidized form of the polymer is electronically conductive, and so a purely electronic (i.e., nonredox) interaction occurs when the metal and CP are brought into electrical contact. The result is an alteration of Fermi energy (that energy at which the probability of occupation of the electron state is 1=2) of both metal and CP, just as when two dissimilar metals are brought into contact. Electrons move from the (more active) metal into the CP until the Fermi energies (or electrochemical potential of electrons, m e) in the two phases are equal. In the presence of an aerated electrolyte, the metal and CP then become a galvanic couple, where oxygen reduction (perhaps catalytic as discussed later in the chapter) can occur at the CP surface. Additionally, a redox-based electron transfer from the metal (which becomes oxidized) to the CP (which becomes reduced) might be expected, in which case, the availability of charge (i.e., oxidation capacity) within the bulk polymer
R +
R
−
+m e−
A N H
n
oxidized or p-doped form
−m e− m
+ m A− N H
n
m
neutral or undoped form
FIGURE 15.2 Redox scheme for polypyrrole, showing anion (A) expulsion upon reduction. Typically n ¼ 2–4, symbolizing a positive charge for every 2–4 monomer units.
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will likely be important. These electronic interactions will be influenced by the quality of the electrical contact between the CP and the metal, which in turn will be influenced by any oxide layer that separates the two. Thus, metal surface preparation and the method of depositing the coating should be important. Finally, chemical interactions between polymer or dopant anions and the metal interface are likely to be important. In fact, known corrosion inhibitors can be incorporated as the dopant anion, in which case the CP functions as an inhibitor release coating. Clearly, a number of complex interactions can occur between a CP and a metal, and we are still a long way from understanding how to design and optimize CP coatings for corrosion control. Indeed, controversy continues to exist over whether conducting polymers will ever be a practical approach to corrosion control. In this chapter, we will discuss the current state of knowledge in this area, using examples from our own work as well as from other laboratories. This chapter will not be an exhaustive review of earlier work, but rather will emphasize advances in our understanding of the fundamental issues noted above. The reader is referred to recent reviews for a more complete literature survey [9,10]. In the remainder of this chapter, we will provide a brief introduction to aqueous corrosion and the traditional methods of corrosion control. As the corrosion of a metal is an electrochemical process, we will briefly summarize the (mainly) electrochemical techniques used for studying corrosion. We will then discuss the important issues surrounding the use of conducting polymers for corrosion control, including approaches to form CP coatings and the role of the oxide layer. We will review recent work describing the use of CPs for the corrosion protection of steels, aluminum and its alloys, and other metals, and conclude with a summary and prognosis.
15.2
Aqueous Corrosion and Passivation
The aqueous corrosion of metals requires the presence of an electrolyte at the metal surface, even if only as a thin film of electrolyte. Ions must move (migrate) to maintain charge balance, as oxidation and reduction reactions occur on the metal surface. Typical oxidation reactions include the oxidation of a metal to either a soluble ionic form or an insoluble form as exemplified by the following reactions: M ! Mnþ þ ne
(15:1)
2M þ 3H2 O ! M2 O3 (s) þ 6Hþ þ 6e
(15:2)
A site at which such oxidation reactions occur is called an anode, and the oxidized products that form are usually influenced by pH value and other electrolyte constituents. The reduction reaction depends on environmental conditions, with the following common reactions: O2 þ 2H2 O þ 4e ! 4OH
(15:3)
2Hþ þ 2e ! H2
(15:4)
A site at which such reduction reactions occur is called a cathode, and it is important to note that these reduction reactions result in a local increase in pH value. Redox reactions involving the CP coating may also be involved, as will be demonstrated later in the chapter. A corrosion cell comprises of an anode and a cathode along with the electrolyte, with which they are in contact, with electrons flowing from anode to cathode through the metal substrate and cations migrating from anode to cathode (anions in the opposite direction) through the electrolyte to maintain an overall charge balance. This ion current in the electrolyte can be mapped (by a method described later in this chapter) and permits the identification of local anodes and cathodes and a measurement of local corrosion rates. It is important to note that the rate of corrosion can be reduced by slowing any one (or more) of the three processes identified
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above: the rate of the oxidation reactions, the rate of the reduction reactions, or the rate of ion migration (i.e., conductivity) between redox sites at the metal surface. The most predictable and manageable form of corrosion is uniform corrosion that results when corrosive attack proceeds evenly over the entire surface, with microscopic anodic areas (where metal dissolution occur) and cathodic areas (where hydrogen evolution or oxygen reduction occur) frequently alternating. General thinning of the metal takes place until failure, although disastrous failures are relatively rare as uniform corrosion is relatively easy to measure and predict. For passive metals that form protective oxide layers (passive films), uniform corrosion results when the passive layer breaks down, more or less, completely. Although uniform corrosion results in only a small fraction of industrial corrosion failures, the total tonnage of metal lost (or wasted) is generally regarded as the highest of all forms of corrosion. For metals that form passive films (e.g., stainless steels and aluminum alloys), the corrosion process is often not uniform, but rather occurs at localized sites on the metal surface at positions where the passive layer has been compromised, leading to pitting corrosion. The passive layer is breached as a result of localized chemical damage to or mechanical disruption of the protective oxide film, promoted by factors such as acidity, low dissolved-oxygen concentrations (rendering the passive film less stable), high concentrations of chloride, or the presence of heterogeneities on the metal surface (e.g., nonmetallic inclusions or second-phase constituents). Pitting is more difficult to detect and predict than uniform corrosion, so it is considered to be far more dangerous. Within the formed pits, highly corrosive conditions can develop, leading to very low pH value and high salt (e.g., chloride ion) concentration. As a result, the pit can propagate rapidly through the metal, leading to perforation of the structure or initiation of a crack. Failure can occur unexpectedly with minimal overall metal loss. Crevice corrosion is another form of localized corrosion that occurs in occluded regions where the electrolyte has limited access. Examples of such regions include under rivets and lap joints. The chemistry of crevice corrosion bears some similarity to pitting corrosion in that aggressive electrolyte conditions can develop in the occluded regions with a resultant acceleration of the corrosion rate. When two dissimilar metals are in electrical contact with one another and both are contacting the same electrolyte, one of the metals will preferentially corrode, a process known as galvanic corrosion (also the principle by which certain types of batteries function). The more active metal will corrode, which is the metal having the more negative open-circuit (or corrosion) potential, when immersed all by itself in the electrolyte; the more noble metal (having the more positive open-circuit potential) will support the reduction reactions. The more active metal is, therefore, the anode and corrodes faster than it would all by itself, whereas the other more noble metal becomes the cathode and corrodes slower than it would alone (or maybe not at all). The electrolyte resistance, important in all corrosion processes, may play a particularly influential role in this type of corrosion process. Galvanic corrosion can occur in engineering structures that are fabricated from two or more metals, each metal selected for its special properties. It can also occur at a single metal alloy surface that is compositionally heterogeneous, an example of which are certain aluminum alloys that have copper-rich second-phase particles embedded in a predominantly aluminum matrix [2]. Sacrificial coatings such as Zn-rich coatings for steel [3] and Mg-rich coatings for aluminum alloys [11] are examples where galvanic corrosion is used to advantage, in this case the more active Zn (or Mg) protecting the more noble Fe (or Al). As noted in Section 15.1, a conducting polymer in electrical contact with an active metal is also expected to function as a galvanic couple, and the importance of this concept will be discussed in greater detail later. The corrosion of alloys is more complex than that of pure metals and is far less understood. Both single phase alloys (e.g., a-brass, a Zn–Cu alloy containing less than 37% Zn) and multiphase alloys (e.g., Cu-containing Al alloys) may undergo dealloying, whereby preferential dissolution of a less-noble component of the alloy occurs. Dealloying can proceed in a uniform manner, leading to a weakened porous structure (as with the dezincification of brass), or on a localized scale, leading to the formation of weak, porous microstructures within the material (as in Al or Mg removal from Al2CuMg or S-phase
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particles in Cu-containing Al alloys [12]). In some cases, dealloying may involve dissolution and redeposition of a more noble constituent [12,13]. Pourbaix diagrams are useful for visualizing and understanding the thermodynamics of corrosion [14]. Such diagrams are constructed for a particular metal from thermodynamic relationships (e.g., Nernst equations and equilibrium constant expressions) involving the various metal-containing species that exist over a range of potential and pH values. The diagrams typically display three types of regions or phase fields in a potential versus pH space: immunity, corrosion, and passivity. The elemental form of the metal (e.g., Al) is stable in the immunity region, typically observed at low potentials where corrosion is thermodynamically unfavorable. The corrosion regions correspond to E–pH conditions over which soluble (ionic) forms of the metal are stable (e.g., Al3þ or AlO2) and, for most metals, are observed at the lower and higher pH regions of the diagram. The passivity regions correspond to E–pH conditions (usually in the mid-pH range) under which insoluble oxide or hydroxide species form (e.g., Al2O33H2O). The formation of such oxidation products on the metal surface often results in a protective barrier that greatly reduces the rate of corrosion. Pourbaix diagrams for aqueous corrosion also include lines that indicate the E–pH behavior of the cathodic reactions expected in water (oxygen and hydrogen ion reductions), thus providing a visual indication of the thermodynamic driving force available for corrosion at a particular pH value. Although Pourbaix diagrams are useful for predicting corrosion behavior and are available for many metals [14], they provide no information on the rates of the reactions that convert one phase of the metal to another. We conclude this section with a brief discussion of typical passive behavior exhibited by many important structural metals, including iron, aluminum, titanium, and chromium. The passive layer that forms on such metals is an oxide or hydrated oxide layer of limited ionic conductivity, a corrosion product that adheres to the metal and provides a barrier to further metal dissolution. In some cases, a salt film precipitation may be involved. Figure 15.3 illustrates an Evans diagram for typical thin film active–passive behavior (e.g., Ni or Cr in sulfuric acid). In this discussion and throughout this chapter, we will consider anodic currents to be positive and cathodic currents to be negative (however, when the logarithm of the current is taken, the sign is discarded). The dashed portion of the curve represents the cathodic branch where metal ion reduction occurs. The solid portion of the curve represents the anodic branch along which the metal oxidation occurs. The intersection point of these two branches corresponds to the equilibrium (reversible) potential E(M=Mþ) and the exchange current density i0(M=Mþ) (an þ electron transfer kinetic parameter) of the M=M redox couple. As the potential is increased positive of E(M=Mþ), the current density (hence, the corrosion rate) increases until a critical passivation potential (Epp) or a critical current density (icrit) is reached, at which point the current drops to a (usually) much lower value (ip), signaling the onset of passivation. In the passive region, the current remains low even though a large driving force (positive potential) is applied to the metal. Above some potential Et, the current again increases as the oxide film dissolves uniformly, often accompanied by oxygen evolution (the transpassive region). Thick film passivity (e.g., Al or Mg in water) differs from that illustrated in Figure 15.3 in that the precipitous drop in current at Epp is not observed, but rather the current levels off from the linear active branch directly to a plateau current, ip. See Ref. [15] for a more detailed discussion on the passivity of metals and alloys. There are other corrosion processes that will not be discussed here, including intergranular corrosion, hydrogen damage, stress corrosion cracking, and corrosion fatigue. The reader is referred to the excellent volume edited by Stratmann and Frankel for a more detailed discussion of all corrosion processes [16].
15.3
Traditional Methods of Corrosion Control
In this section we briefly summarize the more traditional methods for controlling corrosion, including the use of barrier coatings, the use of inorganic and organic inhibitors, and the use of anodic and cathodic protection. More complete discussions of these methods can be found in Refs. [16,17].
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Transpassive region
Substrate Potential
Et
Passive region
icrit
Epp ip
M→M++ e− i0 (M/M+) Active region
E(M/M+) M++e−→M
Log Current density
FIGURE 15.3 Schematic Evans diagram for the behavior of an active–passive metal. E(M=Mþ), reversible potential of the couple; i0(M=Mþ), exchange current density; Epp, passivation potential; icrit, critical anodic current density; ip, passive current density; Et, transpassive potential.
15.3.1 Coatings One of the most common approaches to provide corrosion protection for a metal is to coat the metal with an organic polymer that provides an impervious barrier to the corrosive environment. Often a complete coating system is employed, consisting of two (or more) coating layers, each optimized for a specific function. A primer coating is designed for good adhesion to the metal substrate and may contain corrosion inhibitors, but primer coatings are generally not good barrier coatings. Therefore, a topcoat (a barrier coating applied in one or more layers on top of the primer) is used, often containing pigments that impart aesthetic or other desirable properties to the coating, such as weathering resistance and improved barrier properties. Unfortunately, all barrier coating systems eventually fail, either by ingress of water, oxygen, and ions through natural defects (e.g., pinholes) in the coating or by diffusion of these species through the polymer matrix. Defects introduced accidentally from mechanical abrasion of the coating can also lead to failure. A coating applied to the metal surface also serves to decrease ion mobility at the metal interface, thus increasing the electrolyte resistance between local anodes and cathodes on the metal surface. Metal coatings are also used for corrosion control [18], an example being Zn coatings used on steel (galvanized steel). In this case, the Zn cathodically protects the underlying steel by a sacrificial process (i.e., the Zn is an active coating).
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We note here that a conducting polymer in the p-doped state, whereas neutral in overall charge, contains dopant ions that are either mobile or facilitate ion movement through the polymer by an ion exchange process [19]. Thus, such polymers are not expected to exhibit substantial barrier properties. Coatings based on the undoped form of a conjugated polymer may possess substantial barrier properties, as such polymers are (ideally) free of ions. However, an undoped polymer may undergo spontaneous doping in the presence of oxygen and an electrolyte [20], and any barrier property of the coating may therefore be lost.
15.3.2 Corrosion Inhibitors As barrier coatings ultimately fail or undergo damage, a backup or secondary defense at the damage site against corrosion is usually employed. Chemical corrosion inhibitors are often incorporated into coatings (or added to closed flow systems) and typically function by forming or causing to form an adsorption layer on the metal surface, thereby slowing the rate of either the cathodic reaction (cathodic inhibitors) or the anodic reaction (anodic inhibitors). In some cases, the inhibitor may slow both the anodic and the cathodic reactions (mixed inhibitors) or may form a poor ionically conducting film between the anodic and cathodic sites, thereby increasing the resistance and reducing the corrosion rate (ohmic inhibitors). Some of the most effective inhibitors promote passivation of the metal surface (passivating inhibitors), and may be further classified based on whether the inhibitor is itself an oxidant (oxidizing inhibitors) or requires the presence of oxygen to achieve passivation (nonoxidizing inhibitors). There are over 1100 corrosion inhibitors available for industrial use [21]. A good review of corrosion protection by inhibition has been provided in Ref. [22], and only a few selected examples will be given here. Anodic inhibitors include nonoxidizing passivators (e.g., orthophosphates, molybdates, and benzoates) as well as oxidizing passivators (e.g., chromates, nitrites, and nitrates). Often these inhibitors or their reaction products are incorporated into the passive layer. Cathodic inhibitors often function by film-forming mechanisms, forming surface layers of poor electrical conductivity that restrict the diffusion of oxygen to the metal surface, and include polyphosphates, phosphates, silicates, and lignins. Many of the inhibitors that exhibit mixed-inhibition function either by an adsorption mechanism (amines, amides, aromatic azoles, and sulfur-containing compounds) or by a film-forming mechanism (phosphonates and silicates). Chromate- and phosphate-containing compounds are widely used to form conversion coatings on metal surfaces, thin inorganic layers that form spontaneously when the metal reacts with these species. Conversion coatings are usually applied to improve adhesion of organic coatings as well as to provide improved corrosion protection. Most recently, phosphate conversion coatings are applied primarily to ferrous alloys, whereas chromate conversion coatings are used primarily on light alloys such as aluminum [23]. Chromate conversion coatings (CCCs) and chromate-containing primer coatings are often regarded as ‘‘self healing’’ coatings, as Cr(VI) can be leached from the coating by a contacting solution and can move to coating defects where it functions as a passivating inhibitor. Inorganic inhibitors incorporated into coatings include borate, chromate, molybdate, phosphate, phosphite, or silicate, typically as anticorrosive pigments (salts) of the metals aluminum, barium, calcium, strontium, or zinc. It should be noted that many of the inorganic inhibitors are anions and are, therefore, candidates for incorporation as dopant anions into CP coatings for controlled release. Small organic molecule inhibitors often function by an adsorption mechanism, the extent of adsorption being dependent on a number of variables including the potential of the metal relative to its potential of zero charge (pzc) (for oxide surfaces pzc is dependent on pH value), the structure and the charge (or polarizability) of the inhibitor (adsorbate), the structure of the metal (oxide) surface, and the presence of other species in the electrolyte [22]. Such inhibitors are often of the mixed type, reducing the rate of both cathodic and anodic reactions [22]. It is interesting to note that aniline [24], pyrrole [25], thiophene [26], and the functionalized forms of these molecules exhibit the ability to inhibit
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corrosion, the principle interaction with the metal surface occurring through the heteroatom of the molecule. When conjugated polymers are formed from these monomers, several additional types of interactions with the metal become possible, as will be discussed later.
15.3.3 Cathodic and Anodic Protection Cathodic protection is commonly used to protect buried pipelines and storage tanks, offshore structures and ship hulls, and can be used for virtually all metals. Either an external DC power supply or a sacrificial anode is used to force the potential of the metal to a more negative value such that only cathodic reactions take place at the metal surface to any appreciable extent, greatly reducing the corrosion rate of the metal. When used in combination with a coating, the required DC current is small, i.e., that required to protect metal exposed at coating defects. A sacrificial anode can be used when an electrical power source is not available and is made of a metal that is more active (i.e., more easily corrodible) than the metal to be protected. For example, sacrificial anodes made of aluminum alloys are often used to protect steel ship hulls. The sacrificial metal may be incorporated into a coating, as with the Zn-rich [3] and Mg-rich coatings [11] noted in Section 15.2. Anodic protection is not as widely used as cathodic protection as it is applicable only to active–passive metals. In this case, a DC power supply is used to force the potential of the metal in a more positive direction (to a value in the passive region, Figure 15.3) so as to maintain the metal in the passive state where the corrosion rate is low (just ip ). Anodic protection is a useful corrosion control technique under extremely corrosive environments, such as in strongly alkaline or acidic environments, conditions where cathodic protection is not practical. Another advantage of anodic protection is its low current requirements. In fact, the current applied (ip ) provides a direct means for monitoring the corrosion rate of the system. One disadvantage of anodic protection is the possibility of accelerating corrosion of the metal if proper controls are not implemented and the potential is allowed to become more positive than Et or more negative than Epp (Figure 15.3). Applications of anodic protection include the protection of mild or stainless steel equipment used to handle and store concentrated sulfuric acid, as well as pulp and paper mill digesters and clarifiers and storage tanks. Figure 15.4 illustrates the origin of the corrosion potential and also the principles of cathodic and anodic protection for a single oxidation reaction (M ! Mþ) and a single reduction reaction (Hþ ! H2) occurring at the metal surface (the dashed lines represent the current–potential behavior of the reverse reactions and are not important to the present discussion). Because charge balance must be maintained, the potential is pinned at a value, Ecorr, where the cathodic current and the anodic current are equal (i.e., where the two curves intersect). This corrosion potential (Ecorr) is called a mixed potential, as it is determined by a mixture of two (sometimes more) electrochemical reactions. The anodic current (also the cathodic current, as they are equal) at this potential is the corrosion current (icorr). It is important to note that Ecorr and icorr are influenced by both the thermodynamics of the two reactions, manifested by the equilibrium potentials E(Hþ=H2) and E(M=Mþ), and by the kinetics of the two reactions, manifested by the exchange current densities i0(Hþ=H2) and i0(M=Mþ), and by the slopes of the two linear curves (the Tafel slopes). If the potential of the metal is forced negative of Ecorr, for example to E1 (Figure 15.4), the rate of the metal oxidation reaction (anodic current) is decreased from icorr to i1a and the rate of the reduction reaction (cathodic current) is increased from icorr to i1c. The applied current iappl necessary to accomplish this is the difference between i1c and i1a (Figure 15.4), or as i1a ji1cj, iappl is essentially i1c. Alternatively, anodic protection can provide a similar corrosion rate by forcing the potential of the metal positive of Ecorr and into the passive region, for example to E2 (Figure 15.4), where the corrosion current is ip (note that in this example, ip i1a). The applied current required to maintain the protection is just ip, which is considerably lower than iappl required for cathodic protection. The possibility of using a conducting polymer for the anodic protection of active–passive alloys will be discussed in Section 15.5.
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ip
Substrate potential
E2 i0 (H+/H2) E(H+/H2) H+→H2 Ecorr E1
Ecorr M→M+
iappl
i1a
icorr
i1c
Log Current density
FIGURE 15.4 Schematic Evans diagram illustrating the concepts of corrosion potential, cathodic protection, and anodic protection. E(Hþ=H2), reversible potential for the hydrogen reduction reaction (HRR); i0(Hþ=H2), exchange current density for the HRR on the metal surface; Ecorr and icorr are the corrosion potential and corrosion current, respectively. See text for discussion of the other quantities.
15.4
Techniques for Studying Corrosion
The aqueous corrosion of metals is an electrochemical process, so it is not surprising that electrochemical methods are among the most powerful techniques for studying corrosion and its control by conducting polymers. A detailed discussion of these electrochemical methods is beyond the scope of this chapter and the reader is referred to other sources for a more complete discussion [17,27]. In this section, we provide a brief overview of several useful techniques, preceded by some underlying theory. The methods to be discussed below are subdivided into global techniques, which give a surface-averaged response and local techniques that are capable of providing detailed spatial information.
15.4.1 Background For a single redox reaction occurring at a noble metal electrode (e.g., reaction [Equation 15.4] at a Pt electrode), under conditions such that the current is limited only by the rate of electron transfer, the current density is given by the Butler–Volmer equation [28]: (1 a)nFh anFh i ¼ i0 exp exp ¼ ia ic RT RT
(15:5)
where i0 is the exchange current density, a is the charge transfer coefficient (defined here for the reduction reaction, often 0.5), h is the activation overpotential given by h ¼ E Eeq, n is the
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number of electrons transferred, and R, T, and F have their usual electrochemical meanings. When h (overpotential) ¼ 0, the potential E is the equilibrium potential Eeq (predicted by the Nernst equation) and the net current is zero. At this point, the individual components of the current (ia and ic) are equal and each is equal to i0. As can be seen from Equation 15.5, when h > 0, ia > ic and the net current is anodic (positive), whereas when h < 0, ic > ia and the net current is cathodic (negative). The exchange current io is a measure of kinetic facility and a function of the heterogeneous electron transfer rate constant [28]. When two different redox reactions are occurring on the metal surface (as in Figure 15.4), the current is given by the equation [29] 2:3(E Ecorr ) 2:3(E Ecorr ) exp i ¼ icorr exp ba bc
(15:6)
where icorr and Ecorr are the corrosion current and corrosion potential, respectively, and the b terms (known as the Tafel slopes) are given by ba ¼
2:3RT (1 a)nF
bc ¼
2:3RT anF
(15:7)
each having a value of 0.12 V at 258C for a one electron transfer and for a ¼ 0.5. Equation 15.6 is identical in form to Equation 15.5 and is valid when electron transfer is the only rate limiting process (i.e., ohmic polarization and concentration polarization are absent).
15.4.2 Global Electrochemical Methods As (EEcorr) ! 0 (i.e., for small polarization), the exponential terms in Equation 15.6 may be expanded using a power series, and only the leading order terms retained, leading to the result i ¼ 2:3icorr (E Ecorr )
ba þ jbc j ba jbc j
(15:8)
Differentiation of this expression leads to the Stern–Geary equation [17]: icorr ¼
ba jbc j dE where RP ¼ RP 2:3(ba þ jbc j) di E!Ecorr
(15:9)
RP is called the polarization resistance and can be determined from the slope of the linear portion of the i–E curve in the vicinity of the corrosion potential by using a polarization of only a few mV from Ecorr (this method is known as the polarization resistance method). From the measured RP value and the known (or estimated) b values, the corrosion current can be calculated from Equation 15.9. The corrosion rate is then determined from Faraday’s law [17]: Corrosion rate ¼
icorr a m ¼ nF tA
(15:10)
where a is the atomic weight of the metal, m is the mass lost per unit time t, and A is the area of the metal. Another approach to obtain icorr and Ecorr is to use the Tafel extrapolation method. Here a sufficiently large polarization from Ecorr is used so that one or the other of the terms in Equation 15.6 is negligible. For a positive (or anodic) polarization when (EEcorr) > 0, the second term in Equation 15.6 is negligible (recall that Bc is negative) and the equation may then be rearranged to give an expression relating the anodic polarization (ha) to the anodic current (ia):
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ia ha ¼ E Ecorr ¼ ba log icorr
(15:11)
Similarly, for negative (or cathodic) polarization when (EEcorr) < 0, the first term in Equation 15.6 is negligible and the rearrangement then gives a corresponding expression relating the cathodic polarization (hc) to the cathodic current (ic): hc ¼ E Ecorr
jic j ¼ bc log icorr
(15:12)
Equation 15.11 and Equation 15.12, called Tafel equations, explain the linear relationship between potential and log current observed when only the rate of electron-transfer limits the current (e.g., the active region of Figure 15.3). The Tafel extrapolation method is illustrated in Figure 15.5. The solid curve represents the experimental polarization curve that would be observed, whereas the dashed lines are the extrapolations of the linear branches. The intersection of these extrapolated lines provides the corrosion potential and the corrosion current (Ecorr(1) and icorr(1)). Figure 15.5 also illustrates the behavior when concentration polarization is present and the current is limited, partially or entirely, by the mass transfer of the oxidant (e.g., diffusion of oxygen) to the metal surface. The dotted lines in the figure show limiting currents for two cases, labeled ilim(1) and ilim(2). In the first case, the diffusion limited current ilim(1) is large (perhaps because the concentration of oxygen in solution is large) and therefore does not play a role in limiting
M → M+ Anodic branch slope = ba
Substrate potential
to - ∞ Ecorr(1)
Ecorr(2) icorr(1)
Cathodic branch slope = bc
O2 → OH− ilim(2) = icorr(2)
ilim(1)
Log Current density
FIGURE 15.5 Polarization diagram illustrating the Tafel extrapolation method. The solid curve represents the experimental data, whereas the dashed lines are the extrapolations of the linear branches. The dotted curves illustrate the effect of concentration polarization (illustrated for oxygen diffusion), showing limiting currents for two cases, labeled ilim(1) and ilim(2).
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the corrosion current. The experimental curve would look similar to that in Figure 15.5 (solid curve) except that the cathodic branch would have a limited region of linearity and would follow the dotted line labeled ilim(1). In the second case, the diffusion limited current ilim(2) is smaller and the corrosion rate is limited only by the rate of mass transfer, resulting in a smaller corrosion current (icorr(2) ¼ ilim(2)) and a more negative corrosion potential Ecorr(2). In this case, the experimental curve would be shifted in the negative direction (as though it had slid downward along the extrapolated anodic branch until zero current occurs at Ecorr(2)) and would not exhibit a linear cathodic branch, but would instead transition directly to ilim(2). The Tafel extrapolation method is often used in the study of conducting polymers for corrosion control. In such cases, the redox properties of the polymer must also be included, resulting in a mixed system consisting of at least three processes (e.g., metal oxidation, oxygen or hydrogen ion reduction, and CP oxidation and reduction). This situation will be discussed in Section 15.5. Furthermore, there are a number of situations that can lead to non-Tafel behavior [27], including diffusion limitations (discussed above), ohmic losses, the presence of a buffer in solution, or formation of oxide films. Improper interpretation of the polarization curves for passive systems can lead to highly erroneous conclusions [27]. Cyclic polarization experiments (analogous to cyclic voltammetry) are particularly useful for studying oxide formation [15] and for evaluating localized corrosion in terms of breakdown and repassivation potentials [27]. The preceding polarization methods are often called DC methods in that a DC potential is applied to polarize the metal substrate away from its open-circuit potential (e.g., Ecorr). Except for the polarization resistance method, where a very small polarization is used, the i–E response of these methods is usually nonlinear. Electrochemical impedance spectroscopy (EIS) is an AC method that applies a small (few millivolts) sinusoidal potential perturbation of some frequency f (or v ¼ 2pf ) to the substrate and measures the resulting AC current at that same frequency (alternatively, a current could be applied and the potential measured). The impedance at frequency v is computed from the magnitude and phase angle (with respect to the voltage) of the current response [30]. By repeating the process over a range of frequencies, the impedance spectrum Z(f ) is generated. In corrosion studies, the AC perturbation is usually superimposed on (or added to) the measured open-circuit potential. As the perturbation is small, the electrochemical response is linear as was shown earlier for the polarization resistance method. Consequently, the EIS spectrum is often modeled using circuit elements such as resistors, capacitors, and other elements. Detailed discussions of EIS can be found in Refs. [27,30]. Only a brief overview is provided here. EIS has been used to determine corrosion rates and mechanisms of bare as well as coated metals. Figure 15.6 illustrates the impedance spectrum that might be observed for an electrochemical cell modeled by the equivalent circuit shown in the figure, for example an uncoated corroding metal. This model includes a solution resistance RS, a double layer capacitance Cdl, a polarization resistance RP , and a diffusional impedance ZD. The spectrum of Figure 15.6 is displayed in the Nyquist (or complex plane) format, where the real (or in-phase) part of the impedance Z 0 is plotted on the horizontal or real axis and the imaginary (or out-of-phase) part Z 00 is plotted on the vertical or imaginary axis. As almost all electrochemical cells contain capacitance (e.g., double layer or coating capacitance) and as the impedance of a capacitor (Z ¼ j=vC) lies along the negative imaginary axis, Nyquist plots typically display the negative of the imaginary impedance so that the spectrum lies above the real axis (inductive impedance then results in a portion of the spectrum falling below the real axis). The magnitude of the impedance jZj and the phase angle f can be determined at any frequency f along the curve as illustrated in Figure 15.6. The relationships between jZj, f, Z 0 , and Z 00 are given by jZj ¼
qffiffiffiffiffiffiffiffiffiffiffiffiffiffiffiffiffiffiffiffiffiffiffiffiffiffiffi (Z 0 )2 þ (Z 00 )2
f ¼ arctan
00 Z Z0
(15:13)
Thus, the impedance spectrum can also be displayed by plotting log jZj versus log (f ) and f versus log (f ), called a Bode plot [30]. Both Bode and Nyquist plots are encountered in the literature.
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Rs
Cdl RP
−Z ” (Ω cm2)
ZD f
W ⏐Z⏐ f RS
RS+RP Z ’ (Ω cm2)
FIGURE 15.6 Complex plane (or Nyquist) plot of the impedance spectrum for the equivalent circuit shown. An example impedance vector at some arbitrary frequency is illustrated by the dashed arrow. Frequency increases in the direction shown by the solid curved arrow. Circuit elements: uncompensated solution resistance RS; double layer capacitance Cdl; polarization resistance RP ; and diffusional (Warburg) impedance ZD.
Returning to the equivalent circuit of Figure 15.6, the uncompensated solution resistance RS is obtained as the intercept of the impedance arc (semicircle) on the real axis at high frequency, as the impedance of Cdl becomes zero at high frequency. The diameter of the arc is RP (Equation 15.9) and extrapolation of the arc (dotted line) to the real axis at low frequency gives the sum RS þ RP . At even lower frequencies, the diffusion of some species (perhaps oxygen) contributes to the impedance, indicated by the element ZD. Diffusion is modeled by a Warburg impedance element, essentially a transmission line model. Unbounded (or semi-infinite) diffusion is modeled by an infinite length transmission line and results in the linear ‘‘diffusion tail’’ of slope þ1 observed in the complex plane at low frequency (indicated by W in Figure 15.6). If diffusion is bounded, as in a thin layer of electrolyte or in a thin film of a conducting polymer, then a finite-length transmission model is used with either a short-circuit termination (for a transmissive boundary) or an open-circuit termination (for a blocking boundary) [30]. These finite length Warburg elements result in the diffusion tail bending down toward the real axis (becoming resistive at low frequency) for a transmissive boundary, or bending up to a vertical limit (becoming capacitive at low frequency) for a blocking boundary. The impedance spectrum of a conducting polymer, depending on the film geometry, may display either behavior [31]. A coated metal may be modeled by the equivalent circuits shown in Figure 15.7. Initially, the coating is intact and a single arc in the complex plane is predicted, reflecting the single RC time constant of Figure 15.7a. If the coating resistance is very large, as is often the case, during early stages of exposure, the impedance spectrum will be mostly capacitive and the diameter of the arc very large, so that only a partial arc is observed. As water, oxygen, and ions penetrate the coating, the coating resistance (thus, the diameter of the arc) usually decreases and a second arc may appear, reflecting some areas of delamination of the coating such that electrolyte now contacts the metal surface (Figure 15.7b). Thus, the equivalent circuit models used for corrosion studies often change or evolve over the course of extended exposure. It is often found that the double-layer capacitance or a coating capacitance does not behave like an ideal capacitor, experimentally manifested in the complex plane plot by a depressed semicircle whose center lies below the real axis. This behavior is usually attributed to some distribution (or dispersion) in some physical property of the system (e.g., the porous surface of the metal or the varying thickness or composition of a coating) and is modeled by the use of a constant phase element (CPE) [30].
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Cc Rs
Cdl Rpore
(b) Rp
FIGURE 15.7 Equivalent circuit models for (a) an intact coating and (b) a partially failed coating. Circuit elements: intact coating resistance RC and capacitance CC; coating pore resistance Rpore; other elements as defined in Figure 15.6. RC does not appear in circuit (b) as it is usually so large as to be negligible. In some cases, a diffusional impedance ZD may be needed in series with RP.
EIS data are typically analyzed in terms of an equivalent circuit model by fitting the impedance spectrum calculated from the equivalent circuit model to the experimental data using any of a number of nonlinear least squares fitting packages, such as ZView from Scribner Associates, Inc., Equivcrt by Boukamp [32], or the software available from Gamry Instruments. One caveat should be mentioned: the EIS technique works best for systems that are at equilibrium. It also works for systems that are in steady state, such as corrosion systems, as long as the time required for acquisition of the data is short compared to the time over which the system undergoes significant change. To access the lowest frequencies in EIS, many minutes (or even hours) may be required, and corrosion systems often are not sufficiently stationary to obtain meaningful data at the lowest frequencies. Several examples of the use of Tafel plots and EIS measurements for the study of polyaniline and polypyrrole coatings for corrosion control can be found in the previous edition of this handbook [33].
15.4.3 Local Electrochemical Methods The breakdown of passivity at discrete sites on a metal surface leads to localized corrosion and the flow of ionic currents in the electrolyte between local anodic and cathodic sites. The ionic currents lead to a potential field and associated potential gradients within the electrolyte, the relationship between the local current density and the potential field is given by i (x,y,z) ¼ rV (x,y,z) r
(15:14)
where ¯ı (x, y, z) is the current vector (having magnitude and direction), r is the solution resistivity (V cm) and r V (x, y, z) is the potential gradient (V=cm) at location (x, y, z). By scanning a small reference electrode near the substrate surface in two dimensions and measuring its potential versus a remote reference electrode, a map of the potential variation (but not the potential gradient) near the surface can be obtained. The earliest applications of this scanning reference electrode technique (SRET) used a drawn glass capillary as the scanning microreference electrode [34].
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There are two approaches to measure the local potential gradient, r V (x, y, z). In one approach, a twin-electrode assembly is used consisting of two microreference electrodes (e.g., Ag=AgCl) separated from one another along the direction normal to the sample surface (the z direction) [35]. In the second approach, a single microelectrode is vibrated, either in one direction (the z direction) or in two directions (the x and z directions, each at different vibration frequencies) [36,37]. This second approach, known as the scanning vibrating electrode technique (SVET), possesses a signal-to-noise advantage as lock-in amplifiers are used to measure the voltage amplitudes at the frequencies of vibration. The current density vector can then be determined from Equation 15.14, provided the electrolyte resistivity is known (often obtained from a calibration experiment involving the injection of a known current density from a point source). The vibrating probe is scanned just above the surface of the substrate (which is normally at open circuit), producing a map of DC current density distribution. We have found the SVET to be a useful method, particularly for studying the influence of conducting polymers on the corrosion behavior of various metals [20,38–40], and certain of these results will be summarized later in this chapter. Local electrochemical impedance spectroscopy (LEIS) and local electrochemical impedance mapping (LEIM) have proven useful for studying coated metals, although we are not aware of its application to conducting polymer coatings as yet. These techniques usually employ a three-electrode arrangement to apply an AC potential perturbation to the sample. The local AC current flow through the electrolyte just above the sample surface is then measured, typically using the twin-microelectrode approach [41], although use of a single vibrating microelectrode has also been described [42]. The LEIS experiment involves varying the frequency of the AC perturbation at a fixed position of the twin-electrode (or perhaps at several different fixed positions) above the sample surface. In the LEIM experiment, the twinelectrode probe is scanned in the x–y plane above the sample surface while holding the AC perturbation frequency constant. The use of higher frequencies of the AC perturbation lowers the impedance of dielectric-type samples, such as coated metals. Thus, LEIS and LEIM can be used for both bare and coated metals. This approach has been used to probe intrinsic and extrinsic defects in organic coatings applied to metals [43]. The scanning ion electrode technique (SIET) permits the mapping of pH and metal ion concentrations and is based on a potentiometric measurement at an ion-selective microelectrode [44]. Although this technique has found extensive application for measuring ion gradients in biology, there have been only a few reports of its use in corrosion research [45–47]. The microelectrodes are based either on pulled micropipettes containing liquid-ion exchangers [44] or on solid-state materials, such as Ag=AgCl for chloride measurement [46] or the iridium–iridium oxide electrode for pH measurement [47]. We currently use SIET to map pH changes associated with redox processes occurring at conducting polymer coated metals and an example will be given later in this chapter. We also note that SIET is often referred to as scanning electrochemical microscopy (SECM), but here we will limit the use of the term SECM to experiments that scan a voltammetric (rather than potentiometric) microelectrode above the substrate surface. In SECM, a voltammetric microelectrode or tip (often a Pt disk) of micron or submicron dimension is scanned near the substrate surface, the potential of the tip and (optionally) that of the substrate are independently controlled by a bipotentiostat. The reaction of a redox-active dissolved species at the tip results in a current. As the tip is brought nearer the surface, a change in the tip current is observed, a lower current if the substrate blocks diffusion to the tip (negative feedback), or a higher current if the redox species is regenerated at an electrochemically active surface (positive feedback). The tip current reflects both the tip-substrate distance and the electrochemical properties of the surface. When tip current is plotted as a function of tip position, the resulting three-dimensional image reveals information about the local concentration gradients, electrochemical reactivity of the substrate surface, and surface topography. A number of SECM modes of operation are possible [28,48]. The tip can be used to detect a species generated at the substrate surface (substrate generation=tip collection mode or SG=TC mode), a method that is particularly useful in corrosion studies. Alternatively, the tip generation=substrate collection (TG=SC) mode is useful for kinetic measurements. Additional information may be obtained by
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measuring approach curves (a plot of tip current versus tip-substrate distance) or by scanning the tip potential to generate a voltammogram. SECM has been applied to the study of electron and ion transfer at conducting polymer films [49–51], electron transfer at defect sites in aluminum oxide surfaces [52], cathodic activity at aluminum alloy AA 2024 surfaces [53], corrosion dynamics of stainless steel [54], and for determining local sulfur concentrations dissolved from sulfide inclusions in stainless steels [55]. It is clear that SECM is an important tool for corrosion studies. We are currently using SECM to probe the influence of surface preparation on electron transfer characteristics of aluminum alloy surfaces (important for electrical communication between active coatings and the metal substrate), to the study of electrodeposition of conducting polymers on active metals, and to the determination of local oxygen concentrations above conducting polymer-coated metals. In some cases, we scan a Clark-type oxygen microelectrode just above the substrate surface to map oxygen distribution, a technique that is sometimes called the scanning polarographic electrode technique (SPET).
15.4.4 Surface Spectroscopy and Imaging Methods We conclude this section by mentioning several nonelectrochemical techniques that are useful for investigation of corrosion protection by conducting polymers. X-ray photoelectron spectroscopy (XPS) and Auger electron spectroscopy (AES) are ex situ methods that can provide useful information about the chemical composition of surface and near surface regions of materials. Along with time-offlight and secondary ion mass spectrometry (TOF and SIMS) techniques, these techniques have been used to investigate surface films, oxide layers, and the role of inclusions on pitting corrosion [56]. XPS has also been used to determine the doping level of conducting polymers, an important parameter for understanding the interaction between the CP and an active metal [57]. Samples must be stable under the vacuum conditions of these techniques. X-ray diffraction (XRD) provides information about the phase composition of solid materials and, thus, is useful for the identification of corrosion products. By comparing the diffraction pattern (2u values) with the reference spectra of metal oxides of known composition, modifications of phases that have the same chemical formula can be identified, for example akaganeite, feroxyhyte, and lepidocrocite, all of which have the formula FeO(OH). Scanning electron microscopy (SEM) combined with energy dispersive x-ray analysis (EDXA) provides the morphology and the chemical composition of conducting polymer films and corroding surfaces. An advantage of SEM over optical microscopy is its large depth of field over a wide range of magnification. However, samples must be stable under vacuum conditions and electron irradiation. Conducting coatings (gold or carbon) are often used on insulating samples to prevent electrical charge effects. Both secondary electron images and back-scattered electron images are used, although the latter produces images that enhance the surface topography. The electron beam generates x-rays characteristic of the elements present, and permits elemental mapping of the surface as well as quantitative analysis from micron sized areas. We have used SEM and EDXA to assess CP film morphology [57,58], measure CP film thicknesses [57], and image pit sites beneath CP coatings and perform chemical analysis of such pits [39]. Atomic force microscopy (AFM) and electrochemical atomic force microscopy (ECAFM) have proven useful for the study of nucleation and growth of electrodeposited CP films on Al alloy [59]. AFM was used to study adhesion between polypyrrole and mild steel [60], whereas electric force microscopy (EFM) has been used to study local variations in the surface potential (work function) of CP films [61]. AFM with a conductive tip permits a nanoscale AC impedance measurement of polymer and electrolyte interfaces, permitting differentiation between highly conductive amorphous regions and less-conductive crystalline regions of the CP film [62]. The scanning Kelvin probe (SKP, similar to the EFM technique) permits determination of the work function of CP films. This technique has been used to characterize delamination of a polyaniline containing primer combined with an epoxy topcoat [63] and also to probe the doping-level distribution (from the work function variation) in a conducting poly(2,20 -bithiophene) film [64].
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Conjugated Polymers: Processing and Applications
Conducting Polymers for Corrosion Control—General Considerations
Conducting polymers are able to interact with active metals in several ways, and each type of interaction is capable of altering the corrosion behavior of the metal. In this section, we discuss the variety of ways that a CP coating can potentially influence the corrosion of a metal, recognizing that certain of these influences may be more or less important, depending on the metal, the polymer, the dopant ions, and the environmental exposure conditions. The types of interactions discussed in this section fall into the broad categories of electronic, chemical, and electrochemical, and will be further influenced by the method of preparing the CP coating (discussed in Section 15.6) and also the manner in which the metal surface is prepared (discussed in Section 15.7).
15.5.1 Electronic Interactions Doped conjugated polymers are electronic conductors and, like metals, are characterized by a Fermi level or Fermi energy (Ef ), defined as that energy state for which the probability of occupation is 1=2 [65]. a The Fermi level EFa of a phase a is the electrochemical potential m e of the electron in the material, which a in turn is related to the chemical potential me and the phase potential fa by the expression [28] EFa ¼ mae ¼ mae efa (in eV)
(15:15)
where e is the unit electronic charge. The Fermi energy of conducting polymers is not fixed but can be varied over a range by controlling the doping level of the polymer [66,67]. The doping level in turn can be varied by controlling the voltage applied to the polymer and permitting the current to go to zero. The Fermi energy decreases as the potential (fa) of the polymer is made more positive (Equation 15.15). Differences in Fermi levels can be determined from differences in work functions (defined as the minimum work required to extract an electron into vacuum from the Fermi level of a conducting phase through a surface), which in turn can be measured by the scanning Kelvin probe technique. Oxidized (or doped) forms of conjugated polymers typically have higher phase potentials, lower Fermi energies, and thus higher work functions (>5 eV [66]) than polycrystalline active metals do (e.g., 4.28 eV for Al and 4.5 eV for Fe [68]). When two different conducting phases are brought into electrical contact, electrons flow from the phase having higher Fermi energy to the phase having lower Fermi energy until the Fermi levels in the two phases become equal. If one phase is a p-doped CP and the other an active metal, electrons flow from the metal to CP, producing a Schottky barrier at the interface [69]. For a p-type interface, the Schottky barrier height is the difference between the valence-band maximum of the conducting polymer and the metal Fermi level, although exposure of the interface to water neutralizes the potential drop at the interface due to water dipole orientation in a direction opposite to the intrinsic electric field [69]. When electrolyte contacts the electrically connected pair of conductors, a galvanic couple is then established, resulting in a decrease in phase potential of the CP and an increase in phase potential of the metal (from its original or open-circuit value), leading to the ennobling of CP-coated metals often observed in corrosion studies. This process is illustrated schematically in Figure 15.8 where an active metal and a CP in contact with an electrolyte are electrically connected through a switch. When the switch is open, each of the two phases (metal and CP) has an open-circuit potential (OCP) that reflects the individual Fermi levels of the materials (Equation 15.15). When the switch is closed, the two phases become equal in potential, a galvanic couple is created, and the metal potential is ennobled. We recently reported an experiment in which an Al alloy (AA 2024-T3) was connected via a switch to polypyrrole (PPy) deposited onto indium tin oxide (ITO) glass and the materials immersed in an electrolyte. As the switch was alternately opened and closed, the potentials of the individual materials (switch open) and the coupled materials (switch closed) were measured and the currents above each material were also mapped by the SVET [20]. The
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Switch open • • Eoc Potential i=0 Ecorr Al electrodetr
Ppy film Switch closed • •
Esc
Esc
Potential
i>0 Al electrode
Ppy film
FIGURE 15.8 Schematic diagram of the electronic (galvanic) coupling of an active metal (Al) with a CP film (polypyrrole). The potential scale is with respect to an arbitrary external reference. Ecorr and Eoc denote the opencircuit potentials of the Al alloy and CP film, respectively; Esc denotes the short circuit potential (all are mixed potentials). Note that the potential actually measured for the switch-closed case would depend on the reference electrode placement.
potential measurements are shown in Figure 15.9, where the ennobling of the alloy is clearly evident. As noted in Section 15.4.2, regardless of the switch position, the potentials of these materials are mixed potentials. The ennoblement of the alloy is a consequence of galvanic coupling and should not be taken as an indication of corrosion protection.
15.5.2 Chemical Interactions In this section, we briefly describe purely chemical interactions (not involving formal electron transfer) that can occur between a conducting polymer and an active metal. Dopant ions can also chemically interact at the metal interface and this topic will be considered separately in Section 15.5.6. As conjugated polymers contain p-electron systems and often contain heteroatoms (Figure 15.1), chemical interactions with the metal surface are anticipated. Such interactions are responsible for adhesion of the polymer to the metal surface. Polyaniline is somewhat unique among the conjugated polymers in that it undergoes both electronic and protonic doping reactions (Figure 15.10). The emeraldine base (EB) form of polyaniline is a partially oxidized (50% by definition) form of polyaniline, yet EB has low conductivity and requires protonation (proton doping) to become highly conductive. Some metal ions may also be able to dope polyaniline, a concept that is explored further in Section 15.5.5. The EB form of the polymer has many amine nitrogens that are capable of complexing metal ions. Epstein et al. reported that an EB film led to the extraction of Cu from the top several hundred angstroms of Al 2024-T3 [70], thus removing the Cu=Al galvanic couple that is primarily responsible for the corrosion behavior of this alloy. These complexing abilities of the EB may play a role in this process, although significant electron transfer from a metal to EB has also been observed [71]. Conducting polymers may contain functional groups appended to the polymer backbone, as in the case of poly(2-methoxyaniline-5-sulfonic acid) [72] or poly((4-(3-pyrrolyl))butane sulfonate) [73], and such groups may interact with metal interfaces. Indeed, metal salts were found to cause conformational
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0.2
Potential vs Ag/AgCl (V)
0
Ppy:Switch open
−0.2
Switch closed
−0.4
−0.6
Al alloy:Switch open
−0.8 0
100
200
300
400
500
600
700
Time (min)
FIGURE 15.9 Open-circuit potential of Al alloy (&) and PPy film (~) when the switch was open. Open-circuit potential of the Al alloy coupled to the PPy film (^) when the switch was closed. The switch was alternately opened and closed at each time shown. (Reprinted from He, J., Tallman, D.E., and Bierwagen, G.P., J. Electrochem. Soc., 151, B644, 2004. With permission. Copyright 2004, The Electrochemical Society.)
changes in the water soluble poly(2-methoxyaniline-5-sulfonic acid) from an extended coil to a compact coil structure [72]. The emeraldine salt form of polyaniline (Figure 15.10) may function as a pH buffer at the metal interface, minimizing pH changes that would otherwise occur as a result of redox reactions occurring at the interface. The polyaniline emeraldine salt of p-toluenesulfonic acid (PANI-pTS) was found to effectively inhibit corrosion-driven coating delamination (cathodic disbondment) on an iron substrate [74]. The scanning Kelvin probe was used to study the influence of PANI-pTS volume fraction on delamination rate and on the potential of the intact (undelaminated) coated surface. An inhibition mechanism was proposed in which through-coating cathodic oxygen reduction was suppressed by the ennoblement of the substrate potential and by PANI-pTS-mediated pH buffering of the hydroxide produced by the oxygen reduction reaction.
H N
H H N N +• Polyemeraldine salt (green) −2n H+
+2n H+ H N
H N
N
Polyemeraldine base (blue)
H
+2n H+ +2n e− Reduction
H H N+ N Oxidation 2n H n −2n H+ Polyleucoemeraldine salt (clear) − −2n e +2n H+ −2n H+ +2n H+ +2n e− H H Reduction N N N Oxidation 2n n −2n H+ Polyleucoemeraldine base (clear) − −2n e
N +•
FIGURE 15.10 Polyaniline square scheme showing redox (horizontal) and proton exchange (vertical) reactions. The colors of the films are indicated in parentheses. Only the emeraldine salt form is conductive; anions, not shown, are required for charge balance. Also not shown are the fully oxidized (pernigraniline) forms.
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Finally, we note that doped conjugated polymers can function as ion exchange polymers, and this property has been exploited for the development of sensors, ion exchange membranes, and intelligent microcapsules, the ion exchange property is controlled by the pH or potential [75]. For example, polypyrrole prepared with either chloride or polystyrenesulfonate as dopant ion exhibits preferred anion or cation exchange behavior, respectively [76]. It is not difficult to imagine that these polymers can reduce the concentration or rate of arrival of aggressive ions such as chloride at the metal surface. Indeed, an electropolymerized polypyrrole film containing large-size counterions such as poly(styrenesulfonate) or sodium dodecyl sulfate was found to be permeable only to cations and decreased the progressive invasion of the film by chloride ions [77–79], resulting in a significant improvement in the corrosion protection of iron in a 3% NaCl medium.
15.5.3 Redox Reactions at the Conducting Polymer Surface A conducting polymer can participate in redox reactions in two ways. First, as the polymer is an electronic conductor, electron transfer between redox active species in solution and the CP surface can occur, as observed with metal electrodes. In this case, the CP is often viewed as a mediator that shuttles electrons between the solution species and the underlying conductor that supports the polymer [80]. Second, the CP is itself a redox active material that can undergo oxidation or reduction, a topic that will be discussed in Section 15.5.4. The principal reactions of interest here are the hydrogen-ion reduction reaction (often called the hydrogen evolution reaction, or HER) shown in Equation 15.4 and the oxygen reduction reaction shown in Equation 15.3. The reduction of cationic species at CP films is well known and is sometimes catalytic, with rate constants decreasing with increasing formal potential of the polymer (in qualitative agreement with Marcus theory) [80]. However, reports indicate that HER did not occur on glassy carbon electrodes covered with polypyrrole [80], whereas other workers indicate that this reaction depends on the dopant ion for a Pt electrode coated with poly(N-methylpyrrole) [81]. Polypyrrole, poly(N-methylpyrrole), poly(3-methylthiophene), and polyaniline exhibited catalytic activities toward both oxygen and proton reduction when electropolymerized in the presence of certain Keggin-type heteropolyacid anions [82]. Clearly, the nature of the HER at a conducting polymer surface is an important topic that needs further clarification. Oxygen reduction has been reported to be catalytic at several CPs, including polyaniline, polypyrrole, polythiophene, and poly(3-methylthiophene) films, but not at a poly(3,4-ethylenedioxythiophene) (PEDOT) film [83]. The catalytic activity is attributed to the unique electronic structure of CPs, which leads to chemisorbed oxygen having a fairly high degree of activation (the computed bond length increases by more than 20%), facilitating its reduction at the polymer surface. The reverse reaction (water oxidation) has also been reported to be catalytic at polypyrrole films [84]. The results of a SVET experiment at a PPy-coated Al alloy (AA 2024-T3) are shown in Figure 15.11 [20]. A defect (scribe) was made through the coating to the metal surface. On immersion in dilute Harrison solution (DHS, 0.35% ammonium sulfate, 0.05% sodium chloride), an oxidation current was observed within the defect area and a reduction current was observed to be rather uniformly distributed across the CP surface. From sparging experiments to remove oxygen [20] and from local pH measurements using the SIET (unpublished results), we conclude that most (but not all) of the observed reduction current in Figure 15.11 is due to oxygen reduction at the CP surface. After several hours, the current decreases to near background levels as the defect passivates [20]. The importance of catalytic oxygen reduction at the CP surface is illustrated in Figure 15.12, which shows a schematic Evans diagram illustrating one possible scenario for the influence of the rate of the reduction reaction on corrosion behavior of an active or passive metal. Curve 1 represents slower reduction kinetics as might occur on the metal oxide surface, resulting in a corrosion potential Ecorr(1) and a corrosion current in the active region. More facile reduction kinetics, manifested by larger exchange current density i0 (as suggested above for reduction at a CP film) is illustrated by curve 3, where the corrosion potential is now Ecorr(3) and the corrosion current has been reduced to the passive
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Current density (µA/cm2) −500 500 1000 0
Conjugated Polymers: Processing and Applications
60 0 Y( 0 µm ) −60 0
800 −800
0 ) X (µm
FIGURE 15.11 Current density map (left) and optical micrograph with superimposed current density vectors (right) for a PPy coating on Al 2024-T3 with scribe after 10 min of immersion in 0.35% ammonium sulfate, 0.05% sodium chloride. (Reprinted from He, J., Tallman, D.E., and Bierwagen, G.P., J. Electrochem. Soc., 151, B644, 2004. With permission. Copyright 2004, The Electrochemical Society.)
i02
Substrate potential
i01
i03
Ered Ecorr(3) A
3 Ecorr(1) E(M/M+)
B 2 1
Log Current density
FIGURE 15.12 Schematic Evans diagram illustrating the influence of the rate of the reduction reaction (dotted lines) on active–passive behavior of a metal (solid line). Ered, reversible potential for the reduction reaction; i01, i02, i03, increasing exchange current densities for the reduction reaction; E(M=Mþ), reversible potential for the M=Mþ couple; Ecorr(1) and Ecorr(2) are stable corrosion potentials. Concentration polarization is assumed to be absent.
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current density. Curve 2 illustrates an undesirable situation where two possible (unpredictable) states exist, labeled A and B in Figure 15.12. Other scenarios are possible. For example, for a more active metal, the solid curve of Figure 15.12 would be shifted downward (i.e., E(M=Mþ) would be more negative), such that curve 3 might intersect the solid curve in the transpassive region, possibly accelerating corrosion. Clearly, this is an issue that deserves further study. We conjecture that there may be an advantage in shifting the oxygen reduction reaction away from the metal surface to the CP coating surface and spreading (or diffusing) the reaction rather uniformly across the CP coating surface (as in Figure 15.11), particularly for Cu-containing Al alloys such as AA 2024-T3. At such alloys (either bare or coated with a traditional nonconductive coating), the oxygen reduction reaction typically occurs at the Cu-rich second-phase particles or inclusions, resulting in localized anodic currents in the Cu-depleted regions adjacent to the inclusion, a consequence of increasing ohmic drop through the electrolyte layer at greater distance from the inclusion [2]. This in turn leads to pitting or intergranular corrosion, localized processes than can be exacerbated by differential oxygen concentrations at the metal surface. Additionally, the local pH value increases as a result of the cathodic reaction, possibly leading to a further breakdown of the passive layer or to cathodic disbondment of the coating from the metal surface. By shifting the oxygen reduction reaction to the CP coating and diffusing it across the coating surface, these potential problems may be ameliorated.
15.5.4 Conducting Polymer as an Oxidant As noted in Section 15.1 (introduction to this chapter), conducting (i.e., doped) forms of conjugated polymers are oxidants toward most structural (or engineering) metals. The equilibrium potential of the polymer depends on the monomer used for polymerization, the nature of the dopant ion, the doping level, the electrolyte, and other experimental variables [19]. Placing substituents on the polymer ring system further modifies the equilibrium potential through electronic and steric effects. Table 15.1 shows the standard potentials for several active metals along with the potentials of several oxidants, including common conducting polymers for which the oxidizing strength increases in the order polypyrrole < polyaniline < polythiophene for the parent polymers shown in Figure 15.1. As can be seen from Table 15.1, these polymers have oxidizing power comparable to that of Cr(VI) and oxygen and so are capable of functioning (at least in part) by an anodic protection mechanism. As a result, the doped (i.e.,
TABLE 15.1 Values of Equilibrium Potentials for Various Active Metals and Oxidants Computed at pH 7* Active Metals Mg2þ=Mg Al2O3=Al Al3þ=Al Zn2þ=Zn Fe2þ=Fe Cu2þ=Cu Oxidants Hþ=H2 O2=H2O Cr2O72=Cr3þ Polypyrrole Polyaniline Polythiophene
Potential (V versus Standard hydrogen electrode (SHE)) 2.54 1.93 1.80 0.94 0.62 0.16 Potential (V versus SHE) 0.33 þ0.81 þ0.45 0.1 to þ0.3 þ0.4 to þ1.0 þ0.8 to þ1.2
*Assuming 1 mM concentration for all soluble species except dioxygen that was assumed to be 0.3 mM (air saturated). Source: Tallman, D.E., et al., J. Solid State Electrochem. 6, 73, 2002. Note: For the CPs, a range of potential is provided.
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oxidized) form of conjugated polymers will be most appropriate for protection of metals that exhibit active–passive behavior. In view of the foregoing discussion, there are several important issues that need to be considered in designing and understanding conducting polymers for corrosion control. First, the equilibrium potential of the polymer (e.g., Ered in Figure 15.12) and its relationship to the active–passive behavior of the metal is certain to be important. As noted in Section 15.4.2, the equilibrium potential of the polymer is a mixed potential as other redox reactions such as oxygen reduction can occur at its surface. Thus, predicting the location of the reduction curve (i.e., Ered and i0 in Figure 15.12) and its relationship to the active–passive behavior of the metal is difficult. Polarization experiments are capable of providing this type of information, but most workers fail to recognize that the anodic and cathodic branches of the polarization curve may contain contributions from polymer oxidation and reduction, and, thus, extrapolation of Tafel plots may provide inaccurate estimates of corrosion current. Another issue is the total charge content of the polymer and the portion of that charge that would be available to promote passivation at some defect site. Our group has attempted to address this issue both by calculation of the theoretical charge content of a polypyrrole coating and by experimental measurement of the total charge that can be extracted from the film as well as the charge available through a small opening in a mask (simulating a topcoat with a defect) [20]. From the polymer film density and measured doping level, the total charge content of a 2.7 mm thick film having an area of 1 cm2 was computed to be 145 mC. This theoretical charge was sufficient to account for the current observed in SVET experiments and could sustain a current density of 50 mA=cm2 for 20 h at the 2 2 mm area actually exposed to electrolyte (the remainder of the 1 cm2 area was masked) [20]. Experimental discharge curves were obtained by stepping the potential of PPy-coated ITO from the OCP to 500 mV (versus Ag=AgCl) in nitrogen-sparged electrolyte. When the entire surface of the film was exposed to electrolyte, 55% of the theoretical charge was obtainable after 160 min (the film was still slowly discharging at the end of the experiment, possibly related to slow dopant anion expulsion from interior regions of the film or increased film resistance). When the film was masked such that the area exposed to electrolyte was 250 times smaller than total film area, 18% of the theoretical charge or 33% of the experimental charge was obtainable. Thus, a substantial fraction of the total charge in the bulk of the film is available at a defect and would be capable, for example, of promoting passivation of a metal in a defect region. When the above discharge experiments were repeated in the presence of oxygen, the measured charge increased by a factor of 2.3, verifying that oxygen reduction does take place at the CP film. Furthermore, control experiments at bare ITO and at bare Pt electrodes revealed that the charge from oxygen reduction at these surfaces (under otherwise identical conditions) was only about 1% of that observed at the conducting polymer [20], supporting the earlier suggestion that catalytic oxygen reduction occurs at CP films. To summarize, the conducting polymer by itself can serve as an oxidant, but its behavior is clearly modified by the presence of oxygen, be it by direct catalytic reduction of oxygen at the CP film surface or by oxygen maintaining the polymer in an oxidized state, as discussed in Section 15.5.5. These are important processes by which CP coatings could create and maintain passive films on active metals.
15.5.5 Neutral (Undoped) Forms of Conjugated Polymers Although most studies of corrosion protection by conjugated polymers utilize the doped (thus conducting) form of the polymer, several studies have been conducted on neutral (undoped) polymers. Most of these studies have utilized the EB form of polyaniline and include studies on Al alloys [70,85,86], stainless steel [87], mild steel [88,89], cold rolled steel [90–92], carbon steel [93], and iron [92,94,95]. An undoped poly(2,5-bis(N-methyl-N-hexylamino) phenylene vinylene) or BAMPPV has been studied for the corrosion control of the Al alloy AA 2024-T3 [96,97]. Little else appears to have been published on the use of other undoped conjugated polymers for corrosion control.
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Most of the reports cited above indicate that EB coatings provide some level of corrosion protection on both Fe and Al alloys. The poor adhesion of EB to metal substrates may be a problem [88], and some workers circumvent this problem by incorporating EB particles into a more traditional binder [89,91,98]. As EB is an oxidized form of polyaniline, doping requires only a proton source or perhaps a metal ion source (along with anions to maintain charge balance within the polymer), resulting in conversion of EB to ES. In one study, visual inspection and UV visible spectroscopy indicated that an EB coating on steel was converted to the salt by locally high concentrations of dissolved iron [93]. Another study suggests that electrons are transferred from Fe to EB even under vacuum conditions [71], consistent with a report that EB films (as well as ES films) were able to maintain the potential of stainless steel in the passive region [87]. A chemical interaction between an EB coating on Al and the surface aluminum oxide has been reported, evidenced by changes in the ratio of reduced and oxidized N in the polyaniline as well as changes in the oxide and hydroxide ratio of the aluminum oxide [85]. Thus, it is clear that EB films are capable of reacting with metal and metal oxide surfaces and as a result may be converted to more highly doped forms of the polymer. There is also considerable evidence that neutral forms of conjugated polymers can be oxidized (doped) in the presence of oxygen dissolved in an electrolyte [99,100] and such doping may involve initial formation a charge-transfer complex between the oxygen and the conjugated molecule [99]. Formation of such a complex is consistent with the earlier suggestion (Section 15.5.3) that chemisorbed oxygen having a fairly high degree of activation is responsible for the catalytic oxygen reduction at CP films [83]. From the corrosion literature, one report indicates that galvanic coupling of a steel sample to an ES film resulted in reduction of the polymer to the leucoemeraldine form (Figure 15.10), which was subsequently reoxidized by dissolved oxygen to EB [93]. Another report from our own laboratory demonstrates that the potential of a doped polypyrrole film responds to oxygen concentration (it is a mixed potential; curve a of Figure 15.13) and that a dedoped polypyrrole film undergoes oxygen redoping within 1–2 h (curve b of Figure 15.13) [20].
0.200
0.105
b
0.100
0.100
0.000
0.095 a O2
−0.100
0.090
Potential (V)
Potential (V)
a
N2 −0.200
0.085
O2
b
−0.300
0.080 0
1
2
3
4
Time (h)
FIGURE 15.13 Open-circuit potential as a function of time for a polypyrrole film on ITO exposed to N2-sparged and O2-sparged DHS (0.35% ammonium sulfate, 0.05% sodium chloride). Curve a (right axis): as deposited film immersed first in N2-sparged DHS, then transferred to O2-sparged DHS (indicated by the arrow). Curve b (left axis): A polypyrrole film reduced at 500 mV for 1000 s then immersed first in N2-sparged DHS and then transferred to O2-sparged DHS (indicated by the arrow). (Reprinted from He, J., Tallman, D.E., and Bierwagen, G.P., J. Electrochem. Soc., 151, B644, 2004. With permission. Copyright 2004, The Electrochemical Society.)
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120 100-120 80-100 60-80 40-60 20-40 0-20 −20-0
Current Density (µA/cm2)
100 80 60 40 20 0 388 −20
−269 1069
801 Y (µm)
271
−926 −259
−789
X(µm)
−1805
FIGURE 15.14 Current density map for poly(3-octylpyrrole) on steel after 5 h 35 min immersion in 3% NaCl. The positive (anodic) current was confined to a defect introduced through the coating. (Reprinted from He, J., et al., J. Electrochem. Soc., 147, 3667, 2000. With permission. Copyright 2000, The Electrochemical Society.)
Additional evidence that a nonconductive conjugated polymer film may become conductive on prolonged contact with electrolyte comes from SVET measurements [39]. A poly(3-octylpyrrole) coating was solvent cast onto cold rolled steel and a defect through the coating to the metal surface was introduced. The coating initially was of very low conductivity and no current was detected at the polymer surface; oxidation and reduction currents were confined to the scribe area [39]. After several hours, reduction current rather uniformly distributed over the polymer surface was observed, with only anodic current observed at the defect (Figure 15.14). Similar SVET behavior was observed for an undoped BAMPPV coating having very high initial electrical resistance. Initially, current was confined to the defect area, but after 6 h immersion in electrolyte a uniformly distributed reduction current developed at the polymer surface (unpublished results). Our interpretation of these results is that the CP films were initially nonconductive and unable to support oxygen reduction. On prolonged contact with oxygenated electrolyte, the films became oxidized and sufficiently conductive to mediate the oxygen reduction reaction. We conjecture that in the early stages of exposure, an undoped conjugated polymer film may function as oxygen scavenger, and this may contribute to reported corrosion protection by such films. Eventually, in the presence of metal ions or oxygen, these polymers undergo significant change and the mechanism of interaction with the metal is likely to change. Clearly, studies of undoped conjugated polymers for corrosion control must be conducted, and the results interpreted, with care.
15.5.6 Role of the Dopant Ion The anion that is incorporated into a conjugated polymer on oxidation is called the dopant ion. This anion is released upon reduction of the polymer (Figure 15.2) and may play an important role in the corrosion control mechanism of CP films. Several studies have addressed the influence of the dopant ion on corrosion behavior [40,101–108]. By selecting the dopant ion to be a known corrosion inhibitor, the
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polymer functions as an inhibitor release coating. Such coatings are sometimes referred to as ‘‘smart’’ coatings [102,104,105,109,110] in that the anion is released on demand, i.e., when damage to the coating results in metal oxidation and polymer reduction (galvanic coupling), such as is observed in Figure 15.14. The anion might also be released through an anion exchange process [104]. The released anion can then function to form a second physical barrier to prevent penetration of aggressive ions [109] or can function as an oxygen reduction inhibitor, for example, inhibiting oxygen reduction at Cu-rich cathodic secondary phases in an aluminum alloy [104]. The approach may be particularly valuable for metals where passivity is difficult to achieve, for example, Al alloys in neutral chloride-containing solutions [110]. The oxidative polymerization of polyaniline (PANI) along with 2,5-dimercapto-1,3,4-thiadiazole (DMCT, an anion) led to a mixture of PANI-DMCT and PANI-poly(DMCT). Poly(DMCT) and DMCT form a reversible redox polymerization–depolymerization couple, permitting efficient storage of DMCT in the polymeric form. When this mixture was incorporated into a UV-curable coating system, SVET measurements showed that the coating effectively eliminated corrosion of Al alloy 2024-T3 in pinhole defects [103,105]. DMCT is a well-known oxygen reduction reaction (ORR) inhibitor for copper, and its critical inhibition concentration and mechanism of action has been discussed in Ref. [111]. The corrosion inhibition on the Al alloy was suggested to occur by release of DMCT upon reduction of the polymer mixture, driven by the corrosion process [103]. The DMCT adsorbs onto Cu-rich intermetallic inclusions of the alloy, shutting down the ORR and, thus corrosion. The possibility of continued oxygen reduction on the conducting polymer surface was not addressed. In a similar approach, anionic inhibitors for the ORR, when used as dopants for PANI films, inhibited corrosion of the Al alloy 2024-T3 substrate at a scribe [110]. A PANI formulation doped with phosphonic acid salts and subjected to salt fog exposure revealed the phosphonic acid dopants to be more effective for corrosion protection of mild steel than traditionally used sulfonic-acid salts. SRET data obtained in an acid chloride environment supported the salt fog results, with the sulfonic acid dopants exhibiting an increasing galvanic activity with time, whereas the phosphonic acid dopants showed a decrease in activity with time. It was suggested that passivation of the metal surface occurred through anodization of the metal by the PANI film and formation of an insoluble iron-dopant salt at the metal surface [102]. In another study, a PANI–poly(methyl methacrylate) blend was found to promote corrosion inhibition of steel in chloride solutions by the release of the camphor sulfonate anion dopant [106,109]. It was suggested that such blends have a dual protection mechanism, involving formation of a passivating complex with the dopant anion (camphor sulfonate) and also acting as a physical barrier to chloride anion penetration. It has been found that the surface potential (work function) of conducting polymers as measured by Kelvin probe and by electric force microscopy is influenced by the dopant ion [61], with evidence of local structural inhomogeneity and nonuniform doping-level distribution [64]. For a LiClO4doped poly(2,20 -bithiophene) film, the polymer consisted of grains that varied in work function (and thus in the dopant concentration) from the grain peripheral regions, and the distribution depended on the method of doping [64]. Such spatial variations in surface potential and dopant ion concentration may influence the manner in which conducting polymers interact with active metals at the local level.
15.5.7 Influence of CPs on Overall Coating Impedance We conclude this section by noting an interesting possible consequence of inserting CP material into a coating system. Typically, adding additional layers of coating such that the overall coating thickness is increased results in an increase in the low-frequency (below 0.1 Hz) impedance of the coated sample, a result of increased coating resistance [112]. We have reported that CP films placed between a metal and a topcoat results in lower impedance in the low-frequency region than in control samples where the CP film was omitted [113,114]. Similarly, incorporating 20 wt % polypyrrole into an acrylic paint resulted in a significantly lower impedance than a control sample without polypyrrole, but containing
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20 wt % TiO2 to introduce similar porosity in the coating [115]. The presence of the CP lowered the low-frequency impedance of the coating by approximately two orders of magnitude. Normally, improved corrosion resistance is associated with higher coating impedance, but in the cases cited above, the best corrosion resistance was obtained with coatings exhibiting lower impedance than the respective controls. We have conjectured that the lowering of coating impedance by CPs is because of their ability to conduct current both by ion and electron movement through the polymer [9], thereby facilitating charge transfer between metals (electronic conductors) and electrolytes or barrier coatings (ionic conductors). Ions are readily exchanged at CP and electrolyte or CP and barrier coating interfaces, whereas electrons are readily exchanged at CP and metal interfaces. Thus, overall charge transfer between electrolyte and metal is facilitated and, as a result, the impedance is lowered. This concept is similar to the use of CPs as ion-to-electron transducers in all solid-state ion sensors, where the CP is sandwiched between a metal substrate and an ion-selective membrane [116].
15.6
Approaches to Forming Conjugated Polymer Coatings
Conjugated polymers can be formed by either chemical or electrochemical oxidation of the corresponding monomer, the reaction proceeding through monomeric and oligomeric cation radical intermediates [19]. The polymers thus formed are typically insoluble and nonfusible, making the formation of coatings on active metals difficult. One of the challenges in developing conjugated polymer coatings has been to overcome this difficulty in processing these materials. Several approaches have been used to create coatings made of or containing conjugated polymers. These approaches include adding substituents to the polymer backbone or to the dopant anion that render the polymer solvent soluble or otherwise solvent processible, forming CP–polymer composites or blends (e.g., by incorporating CP particles into traditional polymeric binders), synthesizing copolymers containing CP oligomers, and electrodeposition strategies. We will briefly review these various approaches in this section. We note at the outset that the various approaches to forming CP coatings may result in significant differences in adhesion, conductivity, electroactivity, the amount of CP actually in contact with the metal substrate, and amount of bulk polymer charge available for galvanic coupling.
15.6.1 Functionalized Solvent-Processable Polymers Soluble conducting polymers can be solvent cast to form coatings. The addition of appropriate substituents to the polymer backbone or to the dopant ion can impart the necessary solubility to the polymer. For example, alkyl or alkoxy groups appended to the polymer backbone yield polypyrroles [117,118], polythiophenes [118], polyanilines [119,120], and poly(p-phenylenevinylenes) [97] that are soluble in common organic solvents. Alternatively, the attachment of ionizable functionalities (such as alkyl sulfonates or carboxylates) to the polymer backbone can impart water solubility to the polymer, and this approach has been used to form water-soluble polypyrroles [121], polythiophenes [122], and polyanilines [123]. These latter polymers are often referred to as self-doped polymers as the anionic dopant is covalently attached to the polymer backbone [9]. For use as a corrosion control coating, these water-soluble polymers must be cross-linked [124] or otherwise rendered insoluble. Organic solvent solubility can also be imparted by employing a functionalized dopant ion, often a surfactant-like anion. Pyrrole polymerized in the presence of dodecylbenzene sulfonate resulted in a polymer having solubility in several organic solvents [125,126]. Similarly, a conducting emeraldine salt form of polyaniline was rendered soluble in a range of organic solvents by employing dodecylbenzenesulfonic acid as the dopant ion [127,128]. Colloidal dispersions provide an alternative route for developing solvent-processible conducting polymers and can be produced chemically [129] or electrochemically [130] by the oxidation of monomer in the presence of a steric stabilizer. For electrochemically produced colloids, the steric stabilizer impedes polymer deposition on the electrode surface. Fine colloidal silica can also be used as a dispersant, and nanocomposite colloids of polypyrrole and polyaniline have been prepared using
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chemical oxidation [131] as well as electrohydrodynamic polymerization [132]. In spite of the relatively low percentage of conducting polymer in such composites, they still exhibited reasonable conductivity, and such nanocomposites may have interesting properties when mixed with other polymers or paint formulations for corrosion protection (see Section 15.6.2). Although we refer to certain CPs as soluble, it has been suggested that many soluble CPs are in fact very fine (few nanometers in dimension) colloidal dispersions [133].
15.6.2 Conducting Polymer Composites and Blends Solvent cast coatings such as those described in the previous section often have poor adhesion or cohesion, a consequence of little or no cross-linking within the film. In this section, we discuss approaches that involve blending or polymerizing conducting polymers with other (usually nonconducting) polymers or inorganic fillers to form composite films or coatings. Of course, properties such as conductivity, electroactivity, dopant ion content and availability, and available polymer charge (i.e., oxidation capacity—see Section 15.5.4) will be modified from that of the pure conducting polymer film, and such issues need to be considered. As demonstrated by the following discussion, there has been considerable recent interest in the composite approach. Core-shell particles have been produced both chemically [134] and electrochemically [135]. For example, a dispersion of electrically conductive core-shell particles was obtained by polymerizing pyrrole or aniline in the presence of a dispersion of polyurethane or alkyd resin particles [134]. Coatings from these dispersions were reported to have conductivities in the range of 105 to 10 S=cm [134]. This coreshell approach, though yet to be fully exploited in the area of corrosion protection, enables formation of CP-containing films from waterborne dispersions. A series of conductive electroactive paints were prepared by blending polypyrrole colloids (20–80 wt %) with an acrylic latex formulation in water [136]. Dip coating of lead, stainless steel, and Zincalume substrates resulted in strongly adherent electroactive coatings having electrical conductivities similar to carbon-filled paints. The addition of the CP colloids to the latex paint increased adhesion. Using far less conducting polymer, Wessling has described a polyaniline formulation containing only about 2% PANI, commercially marketed under the trade name CORRPASSIV. Using a dispersion of 70 nm PANI particles, a network of flocculation structures forms at this low PANI concentration, resulting in a conductive primer [137,138]. Heeger has reviewed conducting polymer blends based on PANI and various insulating host polymers, such as polyolefins, poly(methyl methacrylate), polyesters, acrylonitrile butadiene styrene, and poly(vinyl butyral) [139]. The onset of electrical conductivity is observed at volume fractions of PANI below 1%, a consequence of the self-assembled network morphology of the PANI–polymer blends. The blends retain the excellent mechanical properties of the host polymer, and coatings can be fabricated from solution or by melt processing. The PANI network is reported to be sufficiently robust that it remains interconnected and conducting even after removal of the host polymer [139]. The disadvantage of such coatings might be their smaller CP charge content and dopant ion (inhibitor release) capacity. Paints containing higher loadings of PANI have also been described recently [140,141]. An electrically conductive polymer composite of polypyrrole and poly(ethyl methacrylate) has been prepared by an emulsion polymerization procedure [142]. In this case, the relation between conductivity and the polypyrrole content of the composite exhibited a percolation behavior, with conductivities as high as 6–7 S=cm. Such composites might be amenable to melt processing for coating formation. Composite films consisting of polypyrrole or poly(N-ethylaniline) filler dispersed in a polyimide matrix have been described for potential use as corrosion control coatings for the Al alloy AA 2024-T3 [143]. A polyaniline–poly(butyl acrylate-vinyl acetate) composite exhibiting electroactivity and having a conductivity of 2.2 S=cm was prepared by emulsion polymerization. The composite was soluble in common organic solvents and a stable water-based dispersion could also be prepared. Films cast from aqueous media had exceptional mechanical properties and had excellent adhesion to steel [144]. From the same group, a polyaniline and polyvinyl alcohol electroactive composite has been synthesized by
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chemical polymerization of aniline in media containing polyvinyl alcohol (10 wt %). Mechanically robust films could be cast from water-based dispersions of the composite. The conductivity of the films increased with increasing amount of polyaniline in the film to a high value of 2.5 S=cm [145]. Electrosynthesized composites of poly(1,5-diaminonaphthalene) with polyaniline [146] and with polypyrrole [147] have been described and are reported to offer improved corrosion protection of iron. The corrosion-resistant properties of polyaniline and polypyrrole composite coatings electrochemically deposited on low carbon steel [148] were found to be strongly influenced by the applied potential and molar feed ratio of the monomers. The use of polymeric blend composites for corrosion protection of AA 2024-T3 has been reported, including composites formed by incorporating water-soluble conducting polymers (either polymethoxyaniline sulfonic acid or poly(4-(3-pyrrole))butane sulfonate) into various binders (a cross-linked polyvinyl alcohol, a waterborne epoxy, a modified water-dispersible polyester, and a UV-curable urethane acrylate binder) [149]. The preparation of epoxy and polyaniline composite coatings has been described, using either nanodispersed EB particles [91] or EB that was first dissolved in selected amine hardeners before adding the epoxy resin [98]. Even with very low EB loadings, these workers reported enhanced corrosion protection for steel. Lignosulfonic acids are complex polymeric natural products with many random couplings and find commercial use because of their dispersing, binding, complexing, and emulsifying properties. They are commercially produced as a by-product of the paper industry, obtained when processing wood pulp into paper. The polymerization of aniline in the presence of lignosulfonic acid leads to a lignosulfonic acid– doped polyaniline composite (LIGNO–PANI) in which the polymeric dopant is very tightly bound or possibly even grafted to the polyaniline chains [150]. As a result, the composite retains electroactivity even above pH 8. The conductivity of the composite was reported to be between 1 and 5 S=cm. It is suggested that the lignosulfonic acid can be permanently cross-linked into the resin with which it is blended. Cold rolled steel coupons coated with LIGNO-PANI mixed with a water-based acrylic resin were exposed to a 3.55% NaCl solution. Significant corrosion protection was observed for coatings containing as little as 1%–2% LIGNO-PANI. The introduction of Al particles into the compositecontaining resin led to improved performance, possibly a result of cathodic protection of the steel by the Al particles [150]. A number of new polymer composites based on conducting polymers combined with various inorganic materials have recently been described. Polypyrrole and oxide composites have been prepared by the incorporation of Fe3O4 particles into the CP matrix [77,151]. The oxide particles increased the corrosion protection of the films for iron in a 3% NaCl medium, apparently by increasing the oxidizing properties of the film [77] or by helping to maintain the polymer in its oxidized state [151]. A conductive polymer film of polypyrrole doped with polymolybdate anions was electrodeposited onto steel and found to provide corrosion protection in neutral and acidic 3.5% NaCl solution [152]. The anodic codeposition of polypyrrole and TiO2 onto mild steel in an oxalic acid medium has been described [153,154]. The PPy and TiO2 composite showed a considerable improvement in anticorrosion properties with respect to PPy films in salt spray and weight-loss tests. It was suggested that these composite films could be applied as a primary coating replacement for the phosphatized layers on mild steel [154]. A composite polypyrrole coating was formed on an aluminum surface from an aqueous pyrrole solution of fluorozirconic and fluorotitanic acids neutralized with zinc oxide [155]. The composite layer consisted of polypyrrole chemisorbed on titanium and zinc oxides. The coating was reported to exhibit advanced corrosion resistance for Al, with the titanium oxide and zinc playing an important role. Two groups have reported very similar work, namely the use of polyaniline as a conductive matrix for cyanometallate redox centers, such as Prussian Blue or hexacyanoferrates, to form coatings for the protection of stainless steel against pitting corrosion in strongly acid media [156,157]. The coating leads to improved passivation of the steel but also is reported to block access of pitting-causing anions, such as chlorides, to the surface of stainless steel.
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Several recent reports describe using clay or other inorganic fillers to form CP composites. Polyaniline–polypyrrole composite coatings containing clay or yttria stabilized zirconia were electrodeposited onto AA 2024-T3 [158], with improved corrosion resistance of the substrate. Similarly, particulate-filled polyaniline and polypyrrole films on AA 2024-T3 were prepared electrochemically using a variety of fillers, including clay, carbon black, short carbon fiber, zirconia, and silica [159]. Again, enhanced corrosion performance for these composites was observed. Polymer-clay nanocomposites were prepared by intercalating o-ethoxyaniline monomer into the interlayer regions of montmorillonite clay platelets followed by in situ oxidative polymerization, forming the emeraldine base form of the polymer [160]. Coatings from these materials with low clay loadings of up to 3% were prepared on cold rolled steel by drop casting from an NMP solution. Using a similar approach, this same group prepared a dodecylbenzene sulfonate–doped polypyrrole and montmorillonite composite. This composite was electronically conductive, and coatings on CRS were prepared by dip coating from a chloroform solution [161]. In both cases, the composites were found to exhibit much better corrosion protection than the corresponding pristine polymers in 5% NaCl electrolyte. At least some of the improved performance of these filled conducting polymer composites may be attributed to the enhanced barrier property of these materials. Finally, we mention the work of Benicewicz and coworkers who have prepared polymers containing electroactive oligoaniline side chains [162–164]. One might consider these polymers as molecular level composite materials. The flexible backbone structure provides many of the desired physical properties (glass-transition temperature, film-forming properties, etc.), whereas the electroactivity of the material is controlled by the length and number of aniline oligomers in the final polymer. Several (meth) acrylamide and (meth)acrylate monomers containing oligoaniline side-chain units were polymerized via free radical polymerization. The reduced colorless polymer with a trimer side-chain underwent air oxidation to yield a blue color characteristic of the emeraldine base form of polyaniline. When doped with sulfuric acid, the solution turned green characteristic of the emeraldine salt form. The UV–VIS spectroscopy of the polymer was similar to that of polyaniline, and cyclic voltammetry confirmed the electroactivity of the polymers [164]. However, the electrical conductivity of the polymer was low [163], not surprising in view of the lack of extended conjugation in these polymers.
15.6.3 Electrodeposition of Conducting Polymers The electrodeposition of coatings is used extensively in the automotive industry. Such methods are generally considered to be cost effective, readily automated, controllable (film thickness controlled by current, voltage, and time), efficient (high material-transfer efficiency with nearly 100% material utilization and recovery), and environmentally safe (usually a water-based process). The electrodeposition process occurs through local coagulation of a water-dispersed coating due to pH changes at the metal surface resulting from oxidation or reduction of water [165]. Cathodic electrodeposition, where the pH value increases due to water reduction, is used almost exclusively for steels as Fe readily oxidizes during anodic electrodeposition. In contrast, the electrodeposition of conducting polymers is an anodic process that involves both polymerization and subsequent deposition of the polymer. It is most often carried out at noble metal electrodes such as gold or platinum, or sometimes at carbon electrodes, using potentiostatic, potentiodynamic, or galvanostatic methods [19]. These electrochemical methods provide accurate control over the polymerization rate, localizes the polymerization reaction at the metal surface to be covered, and permits precise control of polymer film thickness. The oxidation potentials of active metals are much more negative than those of the monomers that form CPs, and dissolution of the metal occurs before electropolymerization. If polymer does form, it is often only after a lengthy induction period during which the metal dissolution or the passivation occurs. Thus, the direct electrodeposition of CPs onto active metals such as steel or aluminum is complicated by the concomitant oxidation of the metal at the positive potential required for polymerization. In some cases, the formation of an oxide layer on the metal during electropolymerization and deposition may be
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useful in promoting adhesion. In other cases (e.g., Al and its alloys), there is formation of an electrically insulating oxide layer that blocks electron transfer and impedes polymer formation and deposition, leading to patchy, nonuniform polymer films. The characteristics of various metal oxide films and their possible influences on corrosion control by conducting polymers are discussed in more detail in Section 15.7. The hydrogen evolution reaction has also been suggested to interfere with conducting polymer deposition [166]. Here, we summarize recent approaches that have been developed for the direct electrodeposition of CP films onto active metals. Note that many of the CP-composite coatings discussed in the previous section were formed by electrochemical deposition on iron or steel [147,148,151,153,154] or on Al alloy [143,158,159,167]. Work in this area was pioneered by Beck and coworkers [168–172], who focused on the electrodeposition of polypyrrole on iron and aluminum from aqueous and nonaqueous electrolytes. In water, PPy could be electrodeposited onto iron from an electrolyte containing either potassium nitrate or oxalic acid and onto Al from electrolyte containing oxalic acid [170,171]. The oxalate ion is a chelating agent that apparently forms an Fe(II)-oxalate interlayer on iron and produces a porous oxide layer on aluminum [169]. Furthermore, for deposition onto Al, a pretreatment of the aluminum surface by either diamond paste polishing or by anodic activation into the pitting region was required and a high-pyrrole concentration (0.8 M) was necessary. Even under these conditions, the polymerization was still accompanied by growth in the thickness of the Al2O3 layer (Al corrosion) and by the formation of some overoxidized polypyrrole [169]. Overoxidation of conducting polymers is generally undesirable as it increases the localization (i.e., hinders delocalization) of the charge carriers in the polymer, leading to reduced conductivity [19]. Lacroix and coworkers have described the electrodeposition of polyaniline on mild steel by galvanostatic deposition from neutral aqueous solution in the presence of LiClO4, known to passivate mild steel [173,174]. Homogeneous, strongly adherent PANI films were deposited, apparently on the passive layer, with an efficiency of 90% [174]. A two-step process was also described for iron and zinc in which a thin (1 mm) film of PPy was first deposited on the substrate followed by PANI growth. This twostep process apparently involved little or no dissolution of the substrate and the PANI appeared to grow on the surface of the PPy [173]. Adherent and homogeneous PPy films were deposited onto mild steel and zinc from an aqueous medium containing sodium salicylate and pyrrole without the need for metal pretreatment [175,176]. The salicylate forms a passivating, nonblocking layer on a variety of active metals. Current efficiency of the polymerization reaction was close to 100% and the deposition rate was quite high (1 mm=s at 500 mA=cm2). We also note here that this group has reported paint deposition on top of PPy coatings by the cataphoretic technique on mild steel and on galvanized steel, with performance in salt spray tests reported to be as good as that of cataphoretic coatings deposited on phosphated mild steels [177]. Polymethylthiophene (PMT) films were deposited on mild steel after a special pretreatment with 2-(3-thienyl)-ethylphosphonic acid and could be overcoated with conventional topcoats or by cathodic electrodeposition [178]. The PMT film exhibited the ability to protect the substrate exposed through defects in the coating. In another report, electrodeposition of a polythiophene (PTh) film was achieved on PPy-coated mild steel. The initial layer of PPy was deposited by cyclic voltammetry from 0.3 M oxalic acid solution, followed by deposition of the PTh film (also by CV) from an acetonitrile solution of 0.1 M thiophene and 0.15 M LiClO4. EIS, anodic polarization, and open-circuit measurements indicated anodic protection of the mild steel by the PPy=PTh film in 3.5% NaCl [179]. Several ring-substituted anilines (o-toluidine, m-toluidine, o-anisidine, and o-chloroaniline) were electrodeposited on passivated Fe surfaces by cyclic voltammetry, potentiostatic, or galvanostatic techniques from aqueous oxalic acid solutions [180]. With the exception of o-chloroaniline, the films exhibited protective properties against corrosion of Fe in sulfuric acid solution by stabilizing the Fe passive state, though performance was slightly poorer than that of polyaniline. The electrodeposition of PPy films on an aluminum substrate from nitric acid solution was carried out by both potentiodynamic and potentiostatic methods [181]. Oxidation and passivation of the Al occurred first, followed by nucleation and growth of the polymer. It was suggested that polymer
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was initially deposited at flaws in the passive aluminum oxide film by an instantaneous nucleation and tridimensional growth mechanism. More adherent and coherent PPy films were obtained on Al substrates finished with 1000 grit emery paper than for substrates polished with alumina suspensions, attributed to a greater number of flaws and defects in the emery-paper finish. However, the film had poor anticorrosion properties, attributed to the ability of chloride ions to penetrate the film. Three groups have recently reported the successful electrodeposition of PPy onto Cu or Cu alloys. One approach involved either potentiostatic or potentiodynamic electrodeposition on Cu from near-neutral sodium oxalate solution, with film growth apparently facilitated by the initial oxidation of the Cu to form a sufficiently conductive Cu oxalate pseudo-passive layer [182]. This same group reported a similar approach for PPy electrodeposition onto CuZn [183] and CuNi [184] electrodes. For the CuNi electrode, the addition of Cu2þ ions to the electrolyte increased both the rate of the electropolymerization reaction and the adherence of the polymer to the substrate. In the absence of Cu2þ, oxidation of the electrode occurred, generating a Ni-rich layer that was not sufficiently conducting to permit electrodeposition of PPy (possibly due to the formation of NiO with its rather high band gap; see Table 15.2). The PPy films were reported to exhibit significant corrosion protection of both electrodes in acidified and neutral 0.1 M NaCl, even on polarization to high-anodic potentials. A second approach involved the galvanostatic growth of PPy on Cu in the presence of salicylate ions, the coating again reported to provide corrosion protection to the substrate [185]. In a third approach, a PPy film was deposited on brass and on Cu by cyclic voltammetry from a solution containing 0.3 M oxalic acid [186]. Electrochemical impedance spectroscopy (EIS), anodic polarization curves, and open-circuit potential measurements indicated effective corrosion protection of both metals in 0.1 M H2SO4, with somewhat better performance reported for Cu. The electrodeposition of PANI films onto nickel was carried out potentiodynamically, resulting in films with improved stability and adhesion compared to the ones prepared under potentiostatic control [185]. In another study, PPy and PANI films were electrodeposited onto nickel-coated steel panels from TABLE 15.2 Bandgap Energies (Differences between Valence Band and Conduction Band) of Several Metal Oxides for Macroscopically Averaged Films Metal Oxide
Bandgap Energy (eV)
SiO2
9.0
Al2O3 MgO
8.3 7.8
Cr2O3 SnO2 ZnO TiO2 WO3 a-Fe2O3 Au2O3
3.5 3.5 3.2 3.2 2.7 1.9 1.75
PtO PbO2 IrO2
1.3 — —
CuO Cu2O NiO
1.7 1.8 3.5
9 > = > ; 9 > > > > > > > > > =
insulators
> > > > > > > > > ;
n-semiconductors
conductors
9 > =
p-semiconductors
> ;
Source: Schultze, J.W. and Hassel, A.W., Passivity of Metals, Alloys, and Semiconductors, inCorrosion and Oxide Films, eds. M. Stratmann and G.S. Frankel, Wiley–VCH Verlag GmbH & Co. KgaA, Weinheim, Germany, 2003, 216. Note: In the presence of water, the hydrated oxide or hydroxide form is often more stable.
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aqueous solution containing oxalate by cyclic voltammetry [187]. From EIS and anodic polarization studies, the polymer was reported to accelerate and stabilize the formation of Ni(II) oxides, thereby facilitating the repair of any defect caused by a chloride-containing corrosive environment. The PPy coating gave much better performance than did the PANI coating. In many of the studies cited above, electrodeposition of CPs onto oxide-forming metals required (a) a passivating anion to limit dissolution of the metal and (b) a rather high concentration of monomer. Sometimes, a special pretreatment (or polishing) of the metal surface was also required, particularly for metals such as Al that form insulating oxides. The high concentration of monomer (often 0.5 M or greater) may serve to generate a sufficient number of nucleation sites at oxide defects and flaws. If oxidation of the monomer and subsequent formation and deposition of the CP are kinetically limited at a metal oxide surface, then a significant overpotential is required to nucleate and grow the CP film, leading to competitive oxidation processes such as metal oxidation. Our group has been exploring the use of electron transfer mediation for reducing the overpotential required for the electrodeposition of CPs onto active metals, particularly Al and its alloys, thereby alleviating the problems of metal corrosion and polymer overoxidation [57,59,188–191]. A mediator is an electroactive species that thermodynamically is more difficult to oxidize than the monomer to be polymerized, but kinetically it is more readily oxidized (i.e., has lower overpotential) at a metal or metaloxide surface than the monomer. Under such conditions, the mediator is oxidized at the metal substrate (at a lower potential than required for direct monomer oxidation) and then undergoes electron transfer with the monomer, resulting in oxidation of the monomer and regeneration of the reduced form of the mediator (i.e., the process is catalytic). The net result is oxidation of the monomer at a lower potential than would otherwise occur in the absence of the mediator. For example, the compound 4,5-dihydroxy1,3-benzenedisulfonic acid disodium salt (also known as Tiron) has been shown to mediate the electrodeposition of several conducting polymers (including polypyrrole) on platinum electrodes, reducing the deposition potential by up to 200 mV [192]. Figure 15.15 shows the chronopotentiogram for the constant current-(1 mA=cm2) mediated electrodeposition of a PPy film on AA 2024-T3 using Tiron as both mediator and dopant ion. Compared to the nonmediated electrodeposition in the presence of p-toluene sulfonic acid sodium salt (Na-pTS), both the nucleation potential (maximum potential reached in the transient) and the growth (plateau) potential have been lowered by 700 and 500 mV, respectively. The film deposited by Tiron mediation was uniform and complete, whereas the film deposited with Na-pTS was patchy even after two times the deposition time [57]. From measurements of film thickness, doping level and polymer density, the current efficiency for polymer deposition is estimated to be nearly 100%. Electrochemical AFM studies revealed many more nucleation sites during initial stages of electrodeposition in the presence of Tiron than in control experiments where Tiron was replaced by Na-pTS [59]. We note here that the electrodeposition of PPy onto mild steel and zinc in the presence of salicylate [175], described earlier in this section, may in fact have involved electron transfer mediation. The authors noted that salicylate ion was oxidized at platinum and at active metal electrodes and also that pyrrole was electropolymerized at very low potentials in the presence of salicylate. From their voltammograms, we note that PPy was formed at the potential at which salicylate was oxidized, suggesting that electron transfer mediation may have been involved in its electrodeposition (the authors did not consider such a mechanism). More recently, we have reported on the ability of other dihydroxy benzene compounds, including catechol (CAT), hydroquinone (HQ), hydroquinone sulfonate (HQS), and dopamine (DOP), to lower the potential for PPy electrodeposition on AA 2024-T3 and function as electron transfer mediators [191]. We noted that the nonionic mediators CAT and HQ and the cationic mediator DOP (protonated at pH <7) have little influence on the nucleation step, but lowered the potential during the growth stage. On the other hand, the anionic mediators Tiron and HQS lowered both the nucleation potential and the growth potential. At pH < 7, the oxide surface of the alloy is positively charged, the isoelectric point of the various forms of Al2O3 is 7 or higher [193]. Thus, our results suggest that a specific surface interaction of the negatively charged sulfonated mediators with the positively charged oxide surface
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Polypyrrole electrodeposition 2000
Potential (mV vs. Ag/AgCl)
1500
1000 2 500 1 0 Current Density: 1 mA/cm2 −500
1
−1000
2
−1500 −50
Curve 1: with Tiron as counterion Curve 2: with Na-pTS as counterion
150
350
550
750
950
1150
1350
Time (s)
FIGURE 15.15 Potential–time curves for the galvanostatic deposition of polypyrrole on AA 2024-T3 at 1 mA=cm2 current density in the presence of Tiron (curve 1) and in the presence of Na-pTS (curve 2). (Reprinted from Tallman, D.E., et al., J. Electrochem. Soc., 149, C173, 2002. With permission. Copyright 2002. The Electrochemical Society.)
facilitated the nucleation of polymer deposition on the AA 2024-T3 alloy. We further suggested that the anionic mediators (benzene sulfonates) functioned like a surfactant, facilitating access of the rather hydrophobic pyrrole monomer into the rather hydrophilic pores or defects of the oxide surface where mediated electron transfer and, thus, nucleation occur [191]. The mixture of a neutral mediator (CAT) and a nonmediating benzene sulfonate (1,3-benzenedisulfonate) resulted in electrodeposition behavior strikingly similar to that of the anionic mediators, supporting the proposed mechanism. We conclude this section by recalling our assertion made in the introduction to this chapter (Section 15.1) that electronic interactions between a CP film and the underlying metal substrate will be influenced by the quality of the electrical contact between the CP and the metal, which in turn will be influenced by any oxide layer that separates the two. It is reasonable to expect that the electrodeposited CP films will exhibit the highest quality electrical contact (compared for example to solvent cast films), as nucleation and growth of the polymer must occur at surface sites where facile electron transfer through the oxide layer to the underlying metal takes place.
15.7
Oxide Layers and Active Coatings
As corrosion protection by active coatings such as Zn-rich [3], Mg-rich [11], or CP-coatings relies on electrical communication with the underlying metal, the nature of any intervening oxide layer will likely play an important role. As discussed in Section 15.6.3, the electrodeposition of CP films on oxideforming metals is also greatly influenced by the electrical properties of the oxide layer. A detailed understanding of oxide films requires aspects of materials science, solid-state physics, and electrochemistry, and such a discussion is beyond the scope of this chapter. For a detailed discussion of oxide films and their properties, the reader is referred to Ref. [16]. In this section, we provide a brief overview of the
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physical and electronic properties of various metal oxides (with emphasis on those of Fe and Al) and discuss common methods of metal surface preparation and their influence on the oxide layer.
15.7.1 Physical and Electronic Properties of Oxides The rate of electron transfer reactions (ETRs) at a metal surface is strongly influenced by the composition of the surface. As many engineering metals are covered by an oxide layer, the properties of that oxide layer will determine the rate of an ETR (such as oxygen or CP film reduction) or the resistance of an ohmic contact between the metal and a CP film. The band gap energies of several metal oxides are listed in Table 15.2. The conductivity of an oxide film depends on the band gap energy and the position of the Fermi level within the band gap. Foreign ions in the film contribute to defects that serve as donors, acceptors, or traps within the band gap, and modify the conductivity [15]. The thickness of passive films ranges from monolayer for Pt and Au (both of which are low band gap n-type semiconducting oxides) to micrometer dimension for Al (a high band gap, hence insulating, oxide). Thus, rather facile ETR occurs at Pt and Au surfaces, whereas virtually no ETR occurs at a defect-free oxide surface of Al. Real oxide films are typically nonstoichiometric due to an excess of metal ions or a deficiency of oxygen ions in the film and are often amorphous or nanocrystalline. In the presence of water, hydrated oxides or hydroxides often form, such as Al(OH)3 or AlOOH in the passive layer of Al and Fe2O3H2O or g-FeOOH in the passive layer of Fe. Furthermore, the migration or diffusion of defects within the oxide leads to transport of ions within the film and to ion transfer reactions (ITRs) that take place at the oxide–electrolyte interface. Defect concentrations in passive films usually range from 1019 to 1021 cm3 [15]. Thus, as CPs are ion exchange polymers, ion transfer across CP–metal oxide interfaces is likely. Many of the oxides of Table 15.2 are defect semiconductors and most are n-type, a consequence of their tendency to lose oxygen, resulting in an excess of metal (M) atoms capable of donating electrons. Some oxides such as CuO, Cu2O, and NiO are p-type semiconductors [194], a result of their tendency to gain oxygen, resulting in cation vacancies or M3þ sites that act as electron traps or holes. A few oxides have metal-like conductivity, such as PbO2 (important to the functioning of the lead acid battery), whereas others are insulators (e.g., Al2O3). Thus, a wide variation in the rates of ETRs is to be expected at the various oxides, with higher ETR rates at conducting or semiconducting oxides (such as iron or zinc) and low ETR rates at insulating oxides (such as Al). Certain oxides have a rather fixed ratio of cations and anions (e.g., Al2O3), whereas other oxides have a potential dependent composition (e.g., Fe3O4), resulting in a potential dependent conductivity. ETR at oxide-free metals takes place at the Fermi level, but at oxide films ETR takes place at higher- or lower-energy states. The dominating mechanism depends on the band structure of the oxide, oxide film thickness, and the interfacial potential. ETR can take place via the conduction band (n-type oxides) or via the valence band (p-type oxides or n-type oxides at high potential), sometimes involving direct or resonance tunneling [15]. At defect-free high band gap oxides (e.g., Al2O3), ETR cannot occur at all, other than perhaps by tunneling at very thin films. As the majority of studies of CP films for corrosion control are conducted with iron or aluminum alloys, we briefly consider the oxides of these metals in more detail. A large number of iron oxides or hydroxides exist having the basic octahedral unit Fe(O,OH)6, the variation among the different iron oxides and hydroxides are mainly due to a variation in the arrangement of these octahedra. Possible components of the passive film of iron include FeOH2O, Fe2O3H2O, Fe3O4H2O, Fe3O4 (magnetite), aFeOOH (goethite), and g-FeOOH (lepidocrocite). The band gaps of most Fe(III) (hydr)oxide compounds are near 2.0–2.5 eV. However, the bandgap of Fe3O4 (magnetite) is small (0.1 eV), the lowest of any iron oxide [195]. Changes in oxide composition often occur as the sample undergoes dry-wet-dry cycling, leading to an increase in Fe2þ concentration in the oxide layer and the transformation of rather high band gap oxides (such as lepidocrocite) to the near zero band gap magnetite [196]. As a result, a significant increase in the conductivity of the oxide layer occurs during exposure, possibly leading to a change in the mechanism of interaction of the CP film with the metal. Such issues are rarely considered in the literature. One study that has examined the influence of the oxide layer and its
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composition on cold rolled steel and on iron is that by Fahlman et al. [92]. The best corrosion protection of an emeraldine base undercoat was achieved when both the top (a- or g-Fe2O3) and interfacial (Fe3O4) oxide layers were removed before the polymer deposition. An anodic protection mechanism was suggested, with the polyaniline film withdrawing charge from the metal and pacifying its surfaces against corrosion. The oxide layer on Al consists of various solid modifications of hydrated Al2O3, including bayerite, bo¨hmite, corundum, and hydrargillite. The normal surface film formed in air at ambient temperatures is about 5 nm thick and is composed of two layers. The inner-oxide layer next to the metal is a compact amorphous barrier layer, the thickness of which is determined solely by the temperature. At a given temperature, the thickness of this layer is the same in oxygen, in dry air, and in moist air [2]. Covering this barrier layer is a thicker more porous and permeable outer layer of hydrated oxide. At lower temperatures the predominant form of oxide is bayerite, whereas at higher temperatures boehmite is the predominate form. As the oxide film ages, the more stable gibbsite (or hydrargillite, Al2O33H2O) forms and Pourbaix diagrams often show only this most stable form [15]. As noted earlier in this section, ETR at Al oxide films is expected to be blocked by the high band gap of the oxide (Table 15.2). Nevertheless, ETR does occur at Al and Al alloy surfaces, apparently at defects in the oxide layer. Scanning electrochemical microscopy (SECM) was used to study ETR at Al surfaces covered by a 2–3 nm thick native oxide film, using rather pure Al rods and foils [52,197]. Using two different electron transfer mediators (nitrobenzene=nitrobenzene radical anion with E o at 1.6 V versus Ag=Agþ and tetracyanoquinodimethane=tetracyanoquinodimethane radical anion with E o at 0.3 V), similar electrochemical activity was observed, indicating that electrical conduction at the defect sites was only weakly dependent on the interfacial potential and the electric field across the Al2O3 film. The SECM images indicated that the native oxide film on Al contained structural or electronic defect sites of 2–50 mm dimension associated with high electrical conductivity. The density of electroactive defects observed by SECM varied by 2 to 3 orders of magnitude among substrates prepared from the same source of Al, indicating that electrical conduction in the oxide was very sensitive to surface preparation [197]. In another study [53], SECM was performed on Al alloy AA 2024 using the protonated form of (dimethylamino)methylferrocene (DMAFcþ, the Fe2þ form), which was oxidized at the SECM tip to DMAFcþþ (the Fe3þ form). Reduction of DMAFcþþ at cathodically active regions of the substrate led to enhanced current at the tip. Comparison of the SECM images with SEM-energy dispersive spectroscopy images showed that the region’s cathodic activity correlated with the locations of second-phase intermetallic inclusions. The successful electrodeposition of CP films on Al and Al alloys from complexing solutions such as oxalic acid, as discussed in Section 15.6.3, may be attributed to the formation of a porous oxide structure in the presence of such complexing agents [15]. The pores that form have radii of 10–100 nm and depths that can reach tens of microns. The barrier oxide at the base of these pores may be very thin [15] and the nucleation and growth of CP films inside such pores would explain the good adhesion and high electrical connectivity that such films display [57]. Like iron, copper forms duplex oxide layers, i.e., layers with different valency. Both CuO and Cu2O (p-type semiconductors with similar bandgap, Table 15.2) are formed, and higher oxides may exist [15]. The passive behavior of Ni is primarily due to NiO or Ni(OH)2 and NiOOH, although other oxides such as Ni2O3 and NiO2 may also form. Like lead, Ni forms oxides that can be reversibly oxidized and reduced (hence their use in batteries), and such processes will likely need to be considered in understanding how CP films interact with Ni, particularly in view of the rather high bandgap of Ni oxides. To summarize, the ability of an active coating to communicate with the underlying metal substrate will be influenced by any oxide layer that separates the coating from the metal. The type of surface preparation before application of the coating will have a significant influence on the oxide layer and, thus, on the electrical connectivity between coating and metal. For electrodeposited coatings, the method of electrodeposition will also influence the oxide layer. These issues are expected to be particularly important for metals that form high band gap oxides, such as Al and its alloys. Our group is currently studying the influence of surface preparation on electron transfer activity at Al and AA 2024-T3 surfaces.
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15.7.2 Methods of Metal Surface Preparation One of the factors that may be responsible for the sometimes disparate conclusions found in the literature regarding CPs for corrosion control is the wide variation in metal substrate surface preparation. There is very little consistency in the way substrate surfaces are prepared. At the very least, authors should completely describe their surface preparation protocol, including the time interval and conditions of storage (if any) between substrate preparation and coating application. In this section, we briefly review common approaches to substrate surface preparation. Because of the inherent variability of as-received surfaces, the first step is usually the removal of surface films and impurities by abrading and polishing, either by hand or by using a lapping wheel, to generate a fresh and hopefully reproducible surface. The hardness of the abrading material should be approximately three times that of the metal substrate. Common types of abrasives include (increasing in hardness) aluminum oxide, silicon carbide, cubic boron nitride, and synthetic diamond, and these are used either in suspension or in the form of polishing papers. Abrading to a 600 grit finish is usually sufficient to generate a reproducible surface [198] (600 grit corresponds to 16 mm abrading particles). However, the roughness of a freshly prepared surface can influence both the adhesion of a coating and the corrosion rate of the metal. Abrading and polishing are usually performed under water, but such conditions can be aggressive to some alloys. To minimize corrosion during polishing, high-strength Al alloys should be polished in a nonaqueous medium, such as kerosene or absolute ethanol. If polished in aqueous slurries, the reactive intermetallic particles in such alloys will be attacked. Abrading usually begins with a coarse abrasive (<200 grit) and proceeds to a medium abrasive (200–600 grit), sometimes ending with a fine abrasive (>600 grit). Polishing to a mirror finish is usually not practical, particularly for some alloys where differential polishing of the matrix and intermetallic inclusions are problematic. Furthermore, practical application of coatings in the field generally does not involve the use of such fine abrasives. As noted earlier (Section 15.6.3), the adhesion of PPy films on Al was improved for substrates polished with 1000 grit emery paper compared to substrates polished with alumina suspensions [181]. After the grinding and polishing steps, degreasing is typically performed using an organic solvent such as acetone, hexane, or absolute ethanol. The use of acetone to degrease aluminum–copper alloys such as AA 2024-T3 (as per the ASTM E1078-97 cleaning protocol) should probably be avoided, as acetoneinduced pitting of this alloy has been reported, the result of a photocatalytic process occurring at the copper-rich intermetallic particles [199]. Degreasing and deoxidation of metals can also be accomplished using alkaline cleaners, although this approach appears to be rare in conducting polymer studies. Ultrasonic baths are sometimes used in the degreasing process, though the influence of ultrasound on oxide films has not been well studied. One report suggests that cavitation erosion of Al can occur during ultrasonic degreasing [200]. Even if oxide films are completely removed by the above processes, they will spontaneously redevelop at a rate dependent on temperature and humidity. For example, after a few days in humid air, pure aluminum forms thin but dense natural oxide layers 2–10 nm thick, generally considered to consist of amorphous Al2O3. One such air-grown film had a resistance of 35 kV cm2, suggesting that the film structure contained defects or micropores [201]. Thus, it may be important to control or at least specify the environment to which the metal substrate is subjected, if there is a significant time interval between surface preparation and coating application.
15.8
Conducting Polymers for Corrosion Protection—Recent Results
Since our last review of this topic in 2002 [9,10], several new reports on the use of CPs for corrosion control have appeared in the literature. Many of these recent reports have been discussed in previous sections of this chapter. In this section, we summarize new results (from 2002 to the present) not yet discussed in this chapter, focusing on results published in the refereed literature that address corrosion protection mechanisms. The discussion will be organized into three categories: corrosion
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protection of iron and steel, corrosion protection of aluminum and its alloys, and corrosion protection of other metals.
15.8.1 Corrosion Protection of Iron and Steel There have been several new reports of corrosion studies of polyaniline or polyaniline composites on iron or steel [202–214]. In one study, PANI and PPy films were electropolymerized onto iron from aqueous oxalic acid and phosphoric acid solutions, respectively, and copolymers of PANI and PPy were formed on aluminum from a tosylic acid solution [202]. Although all the polymers were reported to exhibit good corrosion protection, the polypyrrole–phosphate system exhibited superior corrosion protection of iron, compared to the polyaniline–oxalate system, attributed to the greater stability of the phosphate layer deposited at the iron substrate, again illustrating the importance of the dopant ion. An alkyd resin containing less than 1% PANI was examined for its ability to protect carbon steel against aqueous corrosion. In field tests, in urban and marine environments, as well as in accelerated laboratory tests, the presence of PANI in the alkyd resin improved the corrosion protection of carbon steel and also the degradation resistance of the coating [214]. Corrosion inhibition of chemically and electrochemically synthesized coatings of PANI, poly (o-methoxyaniline), and their copolymers on stainless steel and on aluminum alloy (6061-T6) was evaluated by immersion in 3% NaCl (steel) and 0.1 M NaCl (Al alloy). These authors concluded from polarization studies that protection involved an anodic protection mechanism [206]. Combination of PANI and PPy films were galvanostatically deposited onto both carbon steel and stainless steel. From potentiodynamic polarization measurements, the performance of the multilayered coatings on carbon steel were not sufficiently better than for single PANI coatings, but for stainless steels the multilayered coatings were reported to be more effective at protecting against pitting corrosion, with films consisting of a PANI layer over a PPy layer providing the best results [213]. The mechanism by which PANI–ES films passivate stainless steel surfaces in H2SO4 has been studied by a variety of experimental techniques, including OCP measurements, Auger depth profiling, and SRET [87,215]. The results suggest that the PANI–ES films hold the potential of the underlying stainless steel in the passive region. The PANI-coated surfaces were found to be enriched in Cr after exposure, as also observed for a bare stainless steel surface whose potential is potentiostatically maintained in the passive region. SRET indicated that the PANI–ES film could passivate exposed stainless steel in the pinhole regions [215]. The more recent study was conducted in dilute H2SO4 solutions (1–500 mM), which are less corrosive but that also result in lower protic doping levels of the PANI film [87]. Exposure to 0.1 mM H2SO4 resulted in the undoped PANI–EB films. Results indicated that both the partially protically doped PANI–ES films and the totally undoped PANI–EB films held the potential of the stainless steel substrate in the passive region, implying that the mechanism by which PANI protects the underlying metal surface from corrosion was independent of doping level. Coatings of chemically produced ES (doped with n-dodecylphosphonate) and EB forms of PANI cast from m-cresol on mild steel were studied by potentiodynamic polarization and EIS techniques, along with a commercially available sulfonic acid-doped PANI [203]. The steel coupons were prepared for coating by anodic cleaning in an alkaline degreaser, followed by an acid pickle and a solvent rinse in xylene, with rinsing between each step using distilled water. A polyacrylic resin topcoat was applied and artificial defects were introduced into polymer-coated steel coupons. Coatings of PANI–ES were found to significantly reduce corrosion rates in 0.1 M NaCl and in 0.1 M HCl, whereas the nonconducting PANI–EB provided some degree of protection in the latter medium. The PANI–ES in electrical contact with the mild steel was reduced to its PANI-LB form, but no reoxidation to the PANI–ES was observed. These authors present convincing arguments that anodic protection in the form of passivation is unlikely in acid or near-neutral chloride environments. Rather, the data support the conclusion that PANI–ES coatings protect mild steel in acidic and near neutral environments by the inhibitory properties of the polymer’s dopant species, which was likely ejected when PANI–ES was reduced as a result of the galvanic coupling between the steel and the CP.
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LEIS, SVET, and micro-Raman spectroscopy were used to probe the mechanism of CP protection at an Fe substrate [212,216]. Both PANI and PPy films were studied using disk and disk-in-ring electrodes, the latter geometry permitting coupling currents to be measured between a bare Fe disk and a CP-coated ring. The results suggest that the CP films were able to passivate small areas of bare iron in defects. The local impedance spectrum displayed characteristics of a passive Fe electrode, with capacitive behavior dominating the entire spectrum. Protection by a PPy film was lost before the entire initial charge stored in the film was consumed, attributed to development of an insulating layer between the Fe and the PPy film [212] or an increase in the electrical resistance of the PPy film [216]. A poorly conducting layer of poly(1,5-diamino-naphthalene) deposited between the Fe substrate and the CP film greatly reduced the coupling current (1 nA), but passivation was still observed. There have been several recent reports of corrosion studies involving polypyrrole or polypyrrole composites on iron or steel [217–226]. One study examined the influence of preparation method on the morphology, mechanical properties, and corrosion inhibition of PPy films on steel, concluding that the best mechanical properties (microhardness, Young’s modulus, and elastic recovery) and the best corrosion protection were obtained for coatings electrodeposited at constant current and then thermally treated at 808C for 1 d [219]. Cyclic voltammetry was used to prepare multilayer coatings on mild steel, including thin polyphenol films on PPy [223,225] and polyindole films on PPy [224]. The corrosion performance of each multilayer coating was compared with that of a single PPy coating, with the multilayer coatings exhibiting improved corrosion performance. Of course, any CP coating will exhibit improved corrosion performance with a topcoat, and even thin polyphenol and polyindol coating provide some topcoat barrier. Upon exposure, most CP coatings do not last long without a topcoat, but then neither do most conventional primers. In situ Raman spectroscopy has been used to study the redox changes (i.e., doping level) that occur in PPy films used as protective coatings on iron by comparing changes in the most sensitive bands to those observed in response to potential variations at a Pt electrode [220]. The films, electrodeposited from oxalate solution, retained their full doping level and maintained the iron in the passive state during the most noble period of the OCP (>0 V versus Ag=AgCl), supporting an anodic protection scheme. A second interval where the OCP was 0.15 V resulted in partial dedoping of the PPy film and was considered to be related to the potential at which ferrous oxalate precipitation healed the formed pits at that potential. A third interval corresponded to the ‘‘death’’ of the film when the OCP fell to the bare iron corrosion potential. A nonuniform doping level was observed at this time, where oxidized (doped) regions of the film, far from the blistered areas, could still be found. Thus, the full oxidizing capacity of the film was not utilized, attributed to low oxalate mobility at this potential. The film was able to be reoxidized by dissolved oxygen, but apparently not sufficiently enough to reach passive conditions again. Sodium oxalate crystals were observed to form around the pits, suggesting that dopant ions may play a role in passivation. A galvanic coupling experiment was described whereby polymethylthiophene (PMT), electrodeposited in the presence hexafluorophosphate onto Pt, was coupled by means of a wire and switch with a mild steel electrode [227]. OCP–time curves were measured while alternately opening and closing the switch and EIS experiments were performed on the separated PMT=Pt and steel substrate. From these experiments, it was concluded that although PMT shifted the potential of the steel positively to 0.5 V versus SCE, no evidence for metal passivation was found. While coupled to the steel, the PMT film resistance increased as reduction of the PMT occurred. Additionally, no evidence for reoxidation of the PMT film by dissolved oxygen was observed. On the other hand, there is probably little or no thermodynamic driving force for the oxidation of polythiophene by oxygen (Table 15.1). Nevertheless, the PMT film deposited on an adhesion promoting self-assembled monolayer on steel, provided corrosion protection as evidenced by potentiodynamic measurements.
15.8.2 Corrosion Protection of Aluminum and Its Alloys There have been a few new reports on the use of PANI for corrosion protection of aluminum and its alloys [228–230]. PANI films electrodeposited at pure aluminum from a tosylic acid containing solution
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lost electroactivity, and EIS data appeared to be strongly influenced by the oxidized Al substrate, apparently a result of a galvanic interaction between the polymer and the aluminum substrate, leading to oxidation of the aluminum and reduction of the polymer [228]. The PANI films offered only a slight increase in corrosion resistance. SVET measurements indicated attack beneath the polymer film, and the authors concluded that chloride anions diffused across the polymer to react at the underlying aluminum substrate. Polymer complexes of PANI and a polymeric dopant such as poly(acrylic acid) or poly (methylvinlyether-alt-maleic acid) have been synthesized using what the authors describe as a ‘‘template-guided’’ synthetic method to form aqueous solutions or dispersions of the complex [230]. An advantage of these polymer complexes is that they remain conductive at pH value up to 9, making them potentially useful in marine environments. The PANI complexes were used as additives (1%–5%) in a commercial waterborne epoxy, which was then applied to various Al alloys (AA 2024-T3, AA 7075-T6, and AA 6061). The coatings were found effective in inhibiting corrosion of the Al alloys in air-saturated seawater. Polymer blends of camphorsulfonate-doped PANI (5–20 wt %) and poly(methylmethacrylate) (PMMA) have been examined for corrosion protection of AA 2024-T3 in sulfuric acid solutions [229]. OCP measurements and Raman spectroscopy showed galvanic coupling between the coating and the alloy, resulting in electrochemical reduction of the PANI component of the blend. SECM revealed that the PANI–PMMA blend drastically suppressed hydrogen evolution in a scratch compared to a control coating of PMMA. The suppression could be mimicked by using a potentiostat to force the substrate potential to the same value at which it is poised by the PANI–PMMA film. Furthermore, the camphorsulfonate anion alone had no inhibitory effect on hydrogen evolution in scratches of AA 2024-T3 coated with only PMMA. By suppressing both the hydrogen evolution and the oxygen reduction reactions, the production of OH near the alloy surface would be inhibited, thereby preventing damage to the alloy surface by dissolution of its protective amphoteric oxide film. It was suggested that this mechanism provides protection against a runaway positive feedback process that would severely damage the alloy surface. Several recent reports on the use of PPy films on aluminum and its alloys were discussed in previous sections of this chapter. One report not previously mentioned is very similar to the work of Saidman and Bessone discussed in Section 15.6.3 [181], wherein PPy films were electrodeposited onto Al substrates from HNO3 containing solution [231]. Corrosion behavior was investigated in HNO3 and KCl solutions using EIS. Although the PPy films caused a positive shift in the OCP as reported by others, corrosion of the substrate in chloride-containing solutions was actually accelerated. Both of these reports indicate a failure of nitrate-doped PPy films to provide corrosion protection of Al, perhaps reflecting the importance of dopant ion in the mechanism (see Section 15.5.6).
15.8.3 Corrosion Protection of Other Metals The use of CP films or CP-containing paints for the corrosion protection of Zn [232], Zn-coated steel [233], Cu [234,235], CuNi [184], and CuZn [183] alloys (see Section 15.6.3), and Mg alloy [236] have been described recently. PANI doped with pTS was dispersed in polyvinylbutyral (PVB) coatings and applied to a zinc substrate [232]. When the coating with a penetrative defect was exposed to aqueous chloride electrolyte, the coatings appeared to inhibit (corrosion-driven) cathodic delamination. The scanning Kelvin probe was used to assess the influence of the volume fraction of PANI=PTS (f) in the coating on delamination kinetics and on the potential of the undelaminated coating surface (Eintact). For f 0.2, no delamination was observed at 95% relative humidity over a 48 h exposure period, even though ennoblement of Eintact did not persist longer than 6 h. A zinc oxide layer is developed at the metal-coating interface, attaining a thickness proportional to f. The authors proposed a mechanism whereby the PANI=PTS induced formation of the oxide layer that, in turn, stifled the O2 reduction reaction and reduced hydroxide production. Additionally, the PANI=PTS served as a pH buffer, reducing alkaline dissolution of the oxide and cathodic delamination of the coating. Potentiodynamic, galvanostatic, and potentiostatic modes were used for a single-step electrodeposition of PPy films on zinc and zinc-coated steel from aqueous electrolytes containing various carboxylate
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salts, including citrate, succinate, oxalate, malate, and tartrate [233]. Only tartrate resulted in adherent and homogeneous PPy coatings, with best results obtained by galvanostatic deposition. OCP and DC polarization measurements in NaCl, HCl, and H2SO4 solutions indicated that the PPy films increased the corrosion potential and reduced the corrosion current for the zinc-coated steel substrates. Interestingly, the use of ultrasound during electrodeposition led to improved resistance to corrosion in salt spray tests. It was suggested that ultrasound caused the formation of cavitation bubbles, which collapsed on the substrate surface, promoting agitation of the solution and enhancing mass transfer. Activation of the substrate surface might also be involved. PPy coatings formed using ultrasound were more compact and adherent. Galvanostatic and potentiostatic techniques were used to deposit PPy films on Cu substrates, with galvanostatic deposition of thick films providing the best corrosion protection in chloride environments [234]. Poly(o-anisidine) (POA) coatings were electrodeposited on Cu by cyclic voltammetry from aqueous solution containing sodium oxalate, and their performance as corrosion control coatings in aqueous 3% NaCl was investigated by potentiodynamic polarization [235]. The coatings were characterized by a variety of techniques, including cyclic voltammetry, UV–VIS absorption spectroscopy, Fourier transform IR spectroscopy, SEM, and x-ray diffraction. The POA film consisted of a mixture of pernigraniline base (PB) and emeraldine salt (ES) forms. The potentiodynamic polarization curves revealed that the POA coating increased the corrosion potential and reduced the corrosion rate of copper by approximately 100-fold (relative to bare Cu). Polypyrrole was electrodeposited on AZ91 Mg alloy having the composition of 9% Al, 1% Zn, and the balance Mg [236]. The PPy films were deposited potentiodynamically from aqueous solutions of 0.2 M pyrrole and 0.2 M KOH and the morphology and composition were characterized by SEM and Auger electron spectroscopy, respectively. Before electrodeposition, the Mg alloy was subjected to a pretreatment, apparently involving electroless deposition of Cu or Ni, although details of this important step are lacking. Both film morphology and shape of the polarization curves were dependent on this pretreatment. A yellow–black PPy coating exhibited cauliflower-like or rod-like structure from the Cu or Ni pretreatment, respectively. Although ennoblement was observed, there apparently was little corrosion protection of the Mg alloy by the PPy coating, perhaps not surprising considering the tremendous difference in equilibrium potentials of these materials (Table 15.1). The incorporation of PPy particles in an acrylic paint appears to be a more promising approach for Mg alloys [115].
15.9
Summary and Prognosis
There are numerous ways in which conjugated polymers, conducting or not, can interact with active metals. Very often, multiple types of interactions (or mechanisms) may be at work, depending on the specific polymer (i.e., its chemical composition), the extent of its doping and the identity of the dopant ions, the characteristics of the active metal and the manner in which its surface is prepared, the manner in which the polymer coating is prepared and applied, pH value of the exposure environment, and the presence of other components of the exposure environment. Interactions between CPs and active metals fall into the three broad categories of electronic, chemical, and electrochemical, and the extent of these interactions will be modulated by the manner in which the CP coating is designed and applied, the way in which the metal surface is prepared, and the nature of any separating oxide layer. Thus, it is neither surprising that a consensus has not been reached regarding ‘‘the mechanism’’ by which the conjugated polymers protect metals against corrosion, nor should we expect such a consensus. The primary modes of interaction can vary widely, and designing an active CP coating for corrosion control of a particular metal will require greater knowledge of the fundamental intricacies of the various interactions discussed in this chapter.
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Acknowledgments We thank our students, postdoctoral associates, and visiting faculty and collaborators, past and present, whose research at North Dakota State University has contributed greatly to our understanding of conjugated polymers and their application to corrosion control. The support of our research by the Office of Scientific Research under Grant F49620-02-1-0398 (Major Jennifer Gresham, Program Officer) is gratefully acknowledged.
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159. Zhu, Y., K. Shah, and J.O. Iroh. 2002. Corrosion protection of Al 2024-T3 by particulate filled conducting polymer coatings., In Surface Modification Technologies XV, 143. Proceedings of the International Conference on Surface Modification Technologies, 15th, Indianapolis, IN. 160. Yeh, J.-M., et al. 2002. Enhancement of corrosion protection effect of poly(o-ethoxyaniline) via the formation of poly(o-ethoxyaniline)-clay nanocomposite materials. Polymer 43 (9): 2729. 161. Yeh, J.-M., C.-P. Chin, and S. Chang. 2003. Enhanced corrosion protection coatings prepared from soluble electronically conductive polypyrrole-clay nanocomposite materials. J Appl Polym Sci 88 (14): 3264. 162. Benicewicz, B.C., and R. Chen. 2000. Synthesis and characterization of polymers with oligoaniline side chains. Polym Prepr Am Chem Soc Div Polym Chem 41 (2): 1733. 163. Chen, R., and B.C. Benicewicz. 2003. Synthesis and characterization of polymers with oligoaniline side chains. In Electroactive polymers for corrosion control, eds. P. Zarras, J.D. Stenger-Smith, and Y. Wei, 126. Washington, DC: American Chemical Society. 164. Chen, R., V. Raghunadh, and B.C. Benicewicz. 2004. New electroactive polymers for anti-corrosion coatings. Polym Prepr Am Chem Soc Div Polym Chem 45 (2): 151. 165. Goldschmidt, A., and H.-J. Streitberger. 2003. BASF handbook on basics of coating technology. Hannover, Germany: Vincentz Network. 166. Huerta-Vilca, D., S.R. Moraes, and A. Jesus Motheo. 2005. Aspects of polyaniline electrodeposition on aluminium. J Solid State Electrochem 9 (6): 416. 167. Shah, K., and J.O. Iroh. 2003. Adhesion of electrochemically formed conducting polymer coatings on Al-2024. Surf Eng 20 (1): 53. 168. Beck, F., and P. Huelser. 1990. Electrodeposition of polypyrrole on aluminum from nonaqueous solutions. J Electroanal Chem Interfacial Electrochem 280 (1): 159. 169. Huelser, P., and F. Beck. 1990. Electrodeposition of polypyrrole layers on aluminum from aqueous electrolytes. J Appl Electrochem 20 (4): 596. 170. Beck, F., P. Huelser, and R. Michaelis. 1992. Anodic deposition of polypyrrole on iron, aluminum, and other commodity metals. Bull Electrochem 8 (1): 35. 171. Beck, F., et al. 1994. Filmforming electropolymerization of pyrrole on iron in aqueous oxalic acid. Electrochim Acta 39 (2): 229. 172. Beck, F., V. Haase, and M. Schroetz. 1996. Polyheteroaromatic layers on commodity metals (CIPL)— passivation, corrosion protection, AIP Conference Proceedings 354 (Organic Coatings), 115. 173. Lacroix, J.C., et al. 2000. Aniline electropolymerization on mild steel and zinc in a two-step process. J Electroanal Chem 481 (1): 76. 174. Nguyen, T.D., et al. 1999. Polyaniline electrodeposition from neutral aqueous media: Application to the deposition on oxidizable metals. Synth Met 102 (1–3): 1388. 175. Petitjean, J., et al. 1999. Ultra-fast electropolymerization of pyrrole in aqueous media on oxidizable metals in a one-step process. J Electroanal Chem 478 (1,2): 92. 176. Aeiyach, S., et al. 2003. Conjugated organic polymers for anticorrosion coating. ACS Symposium Series 832 (Conducting Polymers and Polymer Electrolytes), 128. 177. Ferreira, C.A., et al. 1999. Appraisal of the polypyrrole=cataphoretic paint bilayer system as a protective coating for metals. J Appl Electrochem 29 (2): 259. 178. Rammelt, U., P.T. Nguyen, and W. Plieth. 2001. Corrosion protection by films of intrinsically conducting polymers. Proceedings of Electrochemical Society 2001-22 (Corrosion and Corrosion Protection), 604. 179. Tuken, T., B. Yazici, and M. Erbil. 2004. The use of polythiophene for mild steel protection. Prog Org Coat 51 (3): 205. 180. Sazou, D. 2001. Electrodeposition of ring-substituted polyanilines on Fe surfaces from aqueous oxalic acid solutions and corrosion protection of Fe. Synth Met 118 (1–3): 133. 181. Saidman, S.B., and J.B. Bessone. 2002. Electrochemical preparation and characterization of polypyrrole on aluminum in aqueous solution. J Electroanal Chem 521 (1–2): 87.
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182. Fenelon, A.M., and C.B. Breslin. 2002. The electrochemical synthesis of polypyrrole at a copper electrode: Corrosion protection properties. Electrochim Acta 47 (28): 4467. 183. Fenelon, A.M., and C.B. Breslin. 2003. Corrosion protection properties afforded by an in situ electropolymerized polypyrrole layer on CuZn. J Electrochem Soc 150 (11): B540. 184. Fenelon, A.M., and C.B. Breslin. 2003. The electropolymerization of pyrrole at a CuNi electrode: Corrosion protection properties. Corros Sci 45 (12): 2837. 185. Abrantes, L.M., A.C. Cascalheira, and P.J. Marques. 2002. Anticorrosive protection of nickel and copper by polymeric coatings—influence of experimental conditions. Corrosao e Proteccao de Materiais 21 (1): 14. 186. Tuken, T., B. Yazici, and M. Erbil. 2004. The electrochemical synthesis and corrosion performance of polypyrrole on brass and copper. Prog Org Coat 51 (2): 152. 187. Tuken, T.,et al. 2004. Polypyrrole and polyaniline topcoats on nickel coated mild steel. Prog Org Coat 51 (1): 27. 188. Tallman, D.E., et al. 2003. Electron transfer mediated deposition of conducting polymers on active metals. Synth Met 135–136:33. 189. Vang, C., et al. 2002. Electrocatalytic polymerization of polypyrrole on Al 2024-T3 alloy. Polym Prepr Am Chem Soc Div Polym Chem 43 (1): 742. 190. Tallman, D.E., et al. 2004. Electrodeposition of conducting polymers on active metals by electron transfer mediation. Curr Appl Phys 4:137. 191. Levine, K.L., D.E. Tallman, and G.P. Bierwagen. 2005. The mediated electrodeposition of polypyrrole on aluminium alloy. Aust J Chem 58 (4): 294. 192. Zinger, B. 1988. Catalytic electrosynthesis of conducting polymers. J Electroanal Chem 244:115. 193. Parks, G.A. 1965. The isoelectric points of solid oxides, solid hydroxides, and aqueous hydroxo complex systems. Chem Rev 65 (2): 177. 194. Wang, W., et al. 2003. Synthesis of CuO and Cu2O crystalline nanowires using Cu(OH)2 nanowire templates. J Mater Res 18 (12): 2756. 195. Park, J.-C., et al. 1999. A new synthetic route to wustite. Bull Korean Chem Soc 20 (9): 1005. 196. Leygraf, C. 2003. Atmospheric corrosion. In Corrosion and oxide films, eds. M. Stratmann, and G.S. Frankel, 191. Weinheim, Germany: Wiley–VCH Verlag GmbH & Co. KgaA. 197. Serebrennikova, I., S. Lee, and H.S. White. 2002. Visualization and characterization of electroactive defects in the native oxide film on aluminum. Faraday Discussions 121 (Dynamic Electrode Surface), 199. 198. Samuels, L.E. 2003. Metallographic polishing by mechanical methods, 4th ed. 400. Materials Park, OH: ASM International. 199. Chidambaram, D., and G.P. Halada. 2001. Infrared microspectroscopic studies on the pitting of AA2024-T3 induced by acetone degreasing. Surf Interface Anal 31 (11): 1056. 200. Bol’shakov, L.A., et al. 2002. Kinetics and optimization of aluminum surface degreasing by ultrasound. Izv Vyssh Uchebn Zaved Khim Khim Tekhnol 45 (5): 31. 201. Oh, H.-J., K.-W. Jang, and C.-S Chi. 1999. Impedance characteristics of oxide layers on aluminium. Bull Korean Chem Soc 20 (11): 1340. 202. Breslin, C.B., A.M. Fenelon, and K.G. Conroy. 2005. Surface engineering: Corrosion protection using conducting polymers. Mater Des 26 (3): 233. 203. Cook, A., A. Gabriel, and N. Laycock. 2004. On the mechanism of corrosion protection of mild steel with polyaniline. J Electrochem Soc 151 (9): B529. 204. Ding, K., et al. 2002. Polyaniline and polyaniline-thiokol rubber composite coatings for the corrosion protection of mild steel. Mater Chem Phys 76 (2): 137. 205. Fenelon, A.M., and C.B. Breslin. 2004. Polyaniline-coated iron: Studies on the dissolution and electrochemical activity as a function of pH. Surf Coat Technol 190 (2–3): 264. 206. Huerta-Vilca, D., et al. 2004. PAni as prospective replacement of chromium conversion coating in the protection of steels and aluminum alloys. Mol Cryst Liq Cryst 415:229.
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207. Huh, J., E. Oh, and J. Cho. 2003. Electrochemistry and corrosion characteristics of polyaniline dispersion coating for protection of steels. J Korean Electrochem Soc 6 (2): 113. 208. Iroh, J.O., et al. 2003. Electrochemical synthesis: A novel technique for processing multifunctional coatings. Prog Org Coat 47 (3–4): 365. 209. Martyak, N.M., et al. 2002. Corrosion of polyaniline-coated steel in high pH electrolytes. Sci Technol Adv Mater 3 (4): 345. 210. Meneguzzi, A., M.C. Pham, and C.A. Ferreira. 2002. Electrosynthesis of a double layered intrinsically conducting polymer composite on iron substrates. Mol Cryst Liq Cryst 374:583. 211. Moraes, S.R., D. Huerta-Vilca, and A.J. Motheo. 2003. Corrosion protection of stainless steel by polyaniline electrosynthesized from phosphate buffer solutions. Prog Org Coat 48 (1): 28. 212. Nguyen, T.D., et al. 2004. Mechanism for protection of iron corrosion by an intrinsically electronic conducting polymer. J Electroanal Chem 572 (2): 225. 213. Tan, C.K., and D.J. Blackwood. 2002. Corrosion protection by multilayered conducting polymer coatings. Corros Sci 45 (3): 545. 214. Laco, J.I.I., F.C. Villota, and F.L. Mestres. 2005. Corrosion protection of carbon steel with thermoplastic coatings and alkyd resins containing polyaniline as conductive polymer. Prog Org Coat 52 (2): 151. 215. Gasparac, R., and C.R. Martin. 2001. Investigations of the mechanism of corrosion inhibition by polyaniline. Polyaniline-coated stainless steel in sulfuric acid solution. J Electrochem Soc 148 (4): B138. 216. Nguyen, T.D., M. Keddam, and H. Takenouti. 2003. Device to study electrochemistry of iron at a defect of protective coating of electronic conducting polymer. Electrochem SolidState Lett 6 (8): B25. 217. Asan, A., and M. Kabasakaloglu. 2003. Electrochemical and corrosion behaviors of mild steel coated with polypyrrole. Mater Sci (Translation of Fiziko-Khimichna Mekhanika Materialiv) 39 (5): 643. 218. Hammache, H., L. Makhloufi, and B. Saidani. 2003. Corrosion protection of iron by polypyrrole modified by copper using the cementation process. Corros Sci 45 (9): 2031. 219. Herrasti, P., et al. 2004. Electrochemical and mechanical properties of polypyrrole coatings on steel. Electrochim Acta 49 (22–23): 3693. 220. Nguyen, T.L.H., et al. 2004. Raman spectroscopy analysis of polypyrrole films as protective coatings on iron. Synth Met 140 (2–3): 287. 221. Ocon, P., et al. 2005. Corrosion performance of conducting polymer coatings applied on mild steel. Corros Sci 47 (3): 649. 222. Prissanaroon, W., et al. 2002. Surface and electrochemical study of DBSA-doped polypyrrole films grown on stainless steel. Surf Interface Anal 33 (8): 653. 223. Tuken, T., et al. 2004. The corrosion protection of mild steel by polypyrrole=polyphenol multilayer coating. Corrosion Sci 46 (11): 2743. 224. Tuken, T., et al. 2004. The use of polyindole for mild steel protection. Prog Org Coat 50 (4): 273. 225. Tuken, T., B. Yazici, and M. Erbil. 2004. A new multilayer coating for mild steel protection. ProgOrg Coat 50 (2): 115. 226. Wlodarczyk, R., et al. 2004. Protection of carbon steel against corrosion in aggressive medium by surface modification with polypyrrole film containing hexacyanoferrate. In Trends in Electrochemistry and Corrosion at the Beginning of the 21st Century, eds. E. Brillas and P.L. Cabot, 897. Barcelona: Universitat de Barcelona. 227. Rammelt, U., P.T. Nguyen, and W. Plieth. 2003. Corrosion protection by ultrathin films of conducting polymers. Electrochim Acta 48 (9): 1257. 228. Conroy, K.G., and C.B. Breslin. 2003. The electrochemical deposition of polyaniline at pure aluminum: Electrochemical activity and corrosion protection properties. Electrochim Acta 48 (6): 721. 229. Seegmiller, J.C., et al. 2005. Mechanism of action of corrosion protection coating for AA2024-T3 based on poly(aniline)-poly(methylmethacrylate) blend. J Electrochem Soc 152 (2): B45.
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230. Yang, S.C., et al. 2003. Electroactive polymer for corrosion inhibition of aluminum alloys, in Electroactive polymers for corrosion control, eds. P. Zarras, J.D. Stenger-Smith and Y. Wei, 196. Washington, DC: American Chemical Society. 231. Zhang, A., and X. Guo. 2003. Preparation of Al=Al2O3=PPy and its characteristic of impedance in a chloride-contained solution. Fushi Yu Fanghu 24 (10): 428. 232. Williams, G., et al. 2004. Inhibition of corrosion-driven organic coating delamination on zinc by polyaniline. Electrochem Commun 6 (6): 549. 233. Martins, J.I., et al. 2004. Polypyrrole coatings as a treatment for zinc-coated steel surfaces against corrosion. Corros Sci 46 (10): 2361. 234. Cascalheira, A.C., and L.M. Abrantes. 2004. Polypyrrole films for copper corrosion protection. Corrosao e Proteccao de Materiais 23 (3): 6. 235. Patil, S., S.R. Sainkar, and P.P. Patil. 2004. Poly(o-anisidine) coatings on copper: Synthesis, characterization, and evaluation of corrosion protection performance. Appl Surf Sci 225 (1–4): 204. 236. Jiang, Y.F., et al. 2003. Corrosion protection of polypyrrole electrodeposited on AZ91 magnesium alloys in alkaline solutions. Synth Met 139 (2): 335.
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16 Artificial Muscles 16.1 16.2 16.3 16.4 16.5 16.6 16.7 16.8 16.9 16.10 16.11 16.12 16.13
Toribio F. Otero
16.1
16.14 16.15
Introduction: Muscles and Polymeric Actuators........ 16-1 Electrochemical and Electrochemomechanical Properties of Conducting Polymers ............................ 16-5 Volume Changes............................................................ 16-6 Basic Molecular Actuator and Muscle Similarity ....... 16-7 Devices: Monolayers, Bilayers, Triple Layers, Combinations ................................................... 16-8 Linear Displacement Actuators.................................. 16-14 Triple Layer .................................................................. 16-14 Characterization of the Electrochemomechanical Muscles (Working Potential Ranges)......................... 16-15 Sensing and Actuating Capabilities of Those Electrochemical Reactions .......................................... 16-16 Actuating Capabilities of an Artificial Muscle .......... 16-17 Sensing Capabilities of an Artificial Muscle ............. 16-18 Tactile Sensibility......................................................... 16-20 Mechanical Characterization of Electrochemomechanical Muscles.............................. 16-22 New Expectancies ........................................................ 16-23 Products and Companies Based on CP Artificial Muscles ......................................................................... 16-23
Introduction: Muscles and Polymeric Actuators
Natural muscles are elegant natural devices developed through millions of years of biological evolution to transform chemical energy into mechanical energy and heat. The actuation of a natural muscle (Figure 16.1) involves (a) aqueous media, (b) an electric pulse arriving from the brain (the pulse generator) to the muscle through the nervous system, (c) liberation of calcium ions inside the sarcomere, (d) chemical reactions, (e) conformational changes along natural polymeric chains (actin and myosin) with change of the sarcomere volume, and (f) water interchange. This natural motor can generate quite elegant and gentle movements still never reproduced by any human-made motor. Moreover, the actuation of the muscle involves simultaneous sensing processes providing the living being with a perfect consciousness of both the characteristics of the mechanical movements and the physical interactions between the organ moved by the muscle and its environment: they are intelligent devices. Human technology has been trying to reproduce those movements using electromagnetic motors, internal combustion devices, steam engines, or hydraulic devices. Any machine containing those motors
16-1
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Brain: power generator and ionic conductor
Ionic and chemical stimulus Two ways information
Liberation of calcium ions into sarcomere Chemical reaction Conformational changes on biopolymers
FIGURE 16.1 An electric (ionic) pulse arrives from the brain through nerves to the muscle sarcomere where calcium ions are liberated triggering conformational changes in proteins induced by chemical reactions. All the processes are three dimensional. The generator (brain) and the nerves are ionic conductors. (Modified from Otero, T.F., Polymer Sensors and Actuators, ed. D. de Rossi and Y. Osada, Springer-Verlag, Berlin, 2000. With permission.)
produces quite rudimentary movements and a lot of noise (acoustic and electromagnetic), even though very useful for the human development. In this context scientists had been looking for a different technology capable of producing devices closer to natural muscles or natural molecular motors; hence trying to include, at least, electric pulses and polymeric chains. Some pioneering devices were constructed in the 1950s using films of polymeric gels immersed in aqueous solutions [1–4]. A high electric field was applied across the film using two metallic electrodes, one distant from the gel and one beside the gel. An ionic current flows through the gel and promotes a bending movement of the film due to electroosmotic draining of ions and water from one half of the film thickness, which shrinks toward the other half that swells. When the direction of the current is reversed, the swelling and shrinking processes are also reversed, producing a reverse bending movement. Interaction between electric fields, or electric currents, and swelled polymeric gels attracted the interest of only a few scientists up to the end of the 1980s and the beginning of the 1990s [5,6]. Then a fast development of the field took place due to the interest to reproduce commercial piezoelectric or electrostrictive (electromechanical) devices developed with inorganic materials through the last decades, but using the now similar properties—the control of dimensional changes in polymeric materials by electric fields. The beginning of this explosive interest overlaps the discovery of intrinsically conductive polymers and the reverse variation of their volume when submitted to reverse electrochemical oxidation–reduction processes. This controlled volume variation envisages the construction of new electrochemomechanical devices, bearing new and unexpected actuating and sensing possibilities. All of the developed devices based on the interaction between electric fields or electric currents and polymers were named artificial muscle.
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Summarizing the present state of the art, from the personal point of view of this author, the named artificial muscles, the actuation of which involves polymers, electric fields, and electric currents, can be classified in two main areas: .
.
Artificial muscles (actuators) responding mainly to an electric field, E (electromechanical and electrokinetic devices) Artificial muscles responding mainly to an electric charge, Q—or an electrochemical reaction (electrochemomechanical devices)
The electromechanical polymeric muscles working under the influence of the applied electric field (E) can be classified on the basis of the operating electromechanical transduction phenomena: . . . . .
f (E2): electrostrictive actuators f (E): piezoelectric actuators f (E): Ferroelectric actuators f (E): Electrostatic actuators f (E): Electrokinetic actuators (electroosmotic)
Usually, artificial muscle based on electrostrictive, piezoelectric, electrostatic, or ferroelectric materials have been manufactured as a film of the dry polymer, both sides coated with a thin metallic film required to apply the electric field. Electrokinetic artificial muscles [5,6] are constituted by films of polymeric gel (polymer, solvent, and salt) and two electrodes, located as close as possible to the material or coating both on sides, which are required to apply the electric field that drives the electroosmotic process. Any of the actuators described in this paragraph has a triple layer structure: metal–electroactive polymer–metal (Figure 16.2). The presence of water and ions in electrostrictive, piezoelectric, or ferroelectric polymers produces an overlapping of both actuation processes: electromechanical and electrokinetics. The water also generates a problem: the high applied-potentials (from several volts to thousand of volts) in an electrolytic media
(a) Metal + ∆E
V
Movement
Electroactive Electroactivepolymer Polymer
Metal (b) i
∆E
Electrolyte Conducting polymer Nonconducting polymer Conducting polymer
i
Movement
FIGURE 16.2 (a) Three-layer structure of the electromechanical actuators: A polymeric film, whose dimensions change under an electric field (electroactive material), was coated by two very thin metal films (nonelectroactive) that are required to apply the electric field, acting as a capacitor. (b) Three-layer structure for an electrochemomechanical actuator. A passive film of adherent and elastic polymer is covered by two films of conducting polymer, which simultaneously act as electrodes and electroactive materials (dimensions change under current flow) of an electrochemical cell.
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induce water hydrolysis with the generation of oxygen and hydrogen at the anode–polymer and at the cathode–polymer interfaces, respectively, consuming charge and destroying the device. In any of the electromechanical devices, strong electric fields act on the permanent or induced dipoles present along the polymeric chains promoting coulombic interactions, forcing conformational movements on the polymeric chains and concomitant macroscopic changes of volume, which relax in the absence of the electric field. Similar coulombic interactions occur when a solvent and ions are present, giving electrokinetic (electroosmotic and electrophoretic) processes. So, electrostatic and mechanical models applied to polymeric materials are required to model the attained responses. No chemical reaction is required for the actuation of those devices. In this chapter we focus our attention on polymeric muscles based on the electrochemical properties (those responding to an electric charge by generation of an electrochemical reaction) of the conducting polymers (CP): electrochemomechanical actuators. The theoretical treatment of these devices requires the cooperation of the chemical reactions produced by electrical currents (electrochemistry) involving polymers (polymer science) and ions to generate changes of volume and mechanical energy (thermodynamics and mechanics). The presence of chemical reactions and solvents gives the CP-based actuators a much more close similitude to natural muscles than that of the electromechanical actuators. The interaction between the different chemical and physical properties acting on both equilibrium and rate of any electrochemical reaction provides the electrochemomechanical actuators with unique technological possibilities unimaginable for the electromechanical actuators. Some of them will be shown. The construction of the devices requires the availability of materials, i.e., conducting polymers. Thousands of different materials are being synthesized. Any basic conducting polymer (polypyrrole, polythiophene, polyaniline, etc.) produces a large number of different oxidized materials (polymer– counterion solvent), depending on the salt and solvent used, by chemical or electrochemical oxidation. Moreover, from every basic monomer (pyrrole, thiophene . . . ) seven different families of oxidized materials have been described: (a) basic polymers (polymer–counterion), (b) substituted polymers, (c) self-doped polymers, (d) copolymers, (e) hybrid materials (CP–polyoxometallate), (f) polymeric blends (CP–polyelectrolyte or CP–organic macroanion), and (g) composites based on those families of CPs [7]. Every family of CPs can produce, at least theoretically, a different family of artificial muscles having a different actuation mechanism, different mechanical properties, different voltage actuation, and different problems. A lot of experimental and theoretical work is still required to develop all those possibilities. Due to the complex nature of the polymer–ion interactions (i.e., polymer films act as membranes attaining a thermodynamic equilibrium between anions and cations in the polymer and in solution), any real system involves interchange of either anions or cations during an electrochemical process. Usually one of those interchanges prevails for the different families of CPs. Two main basic ideal electrochemical mechanisms are accepted as the origin of the volume variations: 1. Prevailing anion interchange during redox processes giving an increase of volume during oxidation and a decrease of volume during reduction (accepted for families 1, 2, 4, 5, and 7): (PPy)s þ n(Cl )aq þ mH2 O $ [(PPynþ )s (Cl )n (H2 O)m ]gel þ (ne )metal
(16:1)
2. Prevailing cation interchange shrinking the material during the oxidation processes and swelling during reduction (accepted for families 3 and 6): ! [(PPy nþ )(MA )n ]s þ n(Cþ )aq þ (ne )metal (Red ) (PPy )[n(MA )n(Cþ )]s Ð
(16:2)
(Oxid)
where the different subindexes mean ‘‘s’’ for solid and ‘‘aq’’ for aqueous. MA represents any macroscopic anion trapped inside the CP during polymerization, and PPy can be polypyrrole or any CP.
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The important role played by the water interchange in both oxidized and reduced materials is an unsolved point, in particular when prevailing cation interchanges are present. In most of the systems, during oxidation there coexist the entrance of anions and the expulsion of cations. Most of the literature about the great differences between electromechanical and electrochemomechanical actuators is very confusing, because they are considered as similar devices having similar possibilities and problems. These mixing concepts include theoretical and practical developments, probably founded on the existence of a main common problem: how to produce three-dimensional devices of any shape and volume and able to give any mechanical energy. From conducting polymers, most of the well-stated devices work as very thin bending monolayers, bilayers, or triple layers, or as monolayers, tubes, etc., producing longitudinal movements. Nevertheless, up to now nobody has been able to solve the production of a basic element giving a longitudinal movement and including an actuating electrode, actuating counter-electrode, and reference electrode, which when repeated by n times through the space can give three-dimensional devices having any shape and any volume, whose mechanical properties are n times the property of the basic element. The design has to guarantee the electric contacts and the uniform distribution of the electric current by every constitutive element across the three-dimensional muscle. (A parallel problem for electromechanical actuators requires the uniform distribution of the electric field on the different elements of the three-dimensional device.) In this way, we will try to review the basic molecular aspects that play a major role in the different devices developed based on conducting polymers, their electro–chemo and electromechanical characterization, and their unique simultaneous sensing properties. We will use prevailing anion interchange polymers (anodic swelling and cathodic shrinking). All the used electrochemical, polymeric, mechanical, and physical–chemical concepts are also valid for prevailing cation interchange polymers, as they are the attained conclusions.
16.2
Electrochemical and Electrochemomechanical Properties of Conducting Polymers
As stated in a previous chapter, any film of a basic CP can be used as a working electrode in a solution (solvent þ salt). With the help of a metallic counterelectrode, the material can be submitted to oxidation or reduction processes by allowing the required current flow to maintain the imposed potential (guaranteed by the use of a reference electrode). We describe here those actuation processes where the interchange of anions between the basic polymer and the electrolyte prevails over the interchange of cations (anodic swelling and cathodic shrinking). A close correlation can be stated for those actuation processes based on blends of a basic polymer with a macro anion or a polyelectrolyte, like polypyrrole– polyvinyl sulphonate, or for the family of the self-doped polymers where the interchange of cations prevails. Here the cations are pushed into the polymer during electrochemical reduction and are expelled during oxidation (cathodic swelling and anodic shrinking). So, the process for a polypyrrole film in a solution containing Cl anions can be envisaged as stated above by Equation 16.1. The oxidized polymer is a gel and the electrons are transferred from, or toward, the metal in contact with the conducting polymer. The oxidized polymer [(PPynþ)s(Cl)n (H2O)m]gel is a nonstoichiometric compound: the Cl content can be increased (or decreased) under control of anodic (or cathodic) charges (ne in Equation 16.1). The electrochemical equipment allows a continuous, reversible, and infinitesimal control of the oxidized material composition (a unique fact related to electromechanical actuators) by the flow of constant anodic or constant cathodic currents (constant charge per time unit), by reversing the direction of the current flow, or by the flow of infinitesimal charges, respectively. It is well known that very thin films (nanometers) of reduced materials (CPs) show a very low conductivity, which rises through several orders of magnitude (6–14) when the material is oxidized (doped). Most of this change is produced along low oxidation degrees. The discovery of this unprecedented range for a material property under control earned the chemistry Nobel award 2000 for Professors
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MacDiarmid, Heeger, and Shirakawa. As a natural consequence of the availability of new soft and flexible semiconducting materials, reconstructing electronic and microelectronics producing new flexible and polymeric devices on flexible supports has attracted the interest of a crowd of scientists and engineers. We can estimate that about 80% of the scientists working with conducting polymers are involved in this area. In addition to this unprecedented property, Equation 16.1 contains another unparalleled fact. The nonstoichiometric nature of the oxidized material guarantees that the counterion content (Cl, in this case) can be changed from 0 to 30%–60% (w=w), depending on the counterion used, in a continuous and reversible way by electrochemical reactions consuming electric charges [8,9]. Considering that for films thicker than 1 mm it is very difficult to attain a deep reduction of the material due to their simultaneous structural shrinking and closing entrapping counterions, the concentration range can be considered to move, under electrochemical control, along 4–8 (depending on the material thickness) orders of magnitude [10,11]. This means that any property of the materials changing with their composition will also shift in a continuous, reversible, and infinitesimal way under electrochemical control. It has been well stated that the controlled electrochemical transformation of PPy to [(PPynþ)s(Cl)n (H2O)m]gel promotes the change of different material properties: volume, color, stored charge, porosity, and stored chemicals [7,12–14]. In this chapter we focus our attention on the consequences related to the volume variations linked to electrochemical oxidation and reduction processes in CPs.
16.3
Volume Changes
The electrochemically stimulated conformational relaxation model provides a good description of the electrochemical behavior, including volume changes (Figure 16.3) of conducting polymers [15–26]. In a film constituted by lineal ideal chains of an amorphous conducting polymer, strong Van der Waals attractive interactions are present, giving a compact structure containing a minimum free volume between chains. During oxidation in the presence of an aqueous electrolyte, positive charges are generated along the chains. The distribution of the double bonds along the polymeric chains changes (Figure 16.4), as do angles between consecutive monomeric elements, inducing conformational changes. Positive charges on neighboring chains produce repulsing forces and conformational relaxation movements with generation of free volume between chains. Charge balancing counterions, and the solvent, penetrate from the solution to occupy the newly created free volume, changing the composition of the material. So the generated free volume (and the volume increment of the material) is correlated to the positive charges generated on the neighboring chains: this is the electrochemical charge consumed by the electrochemical reaction by the increase of volume on any basic conducting polymer in the presence of small anions
Polymeric chains
Anions
Oxidation n
+
+ n e− Reduction Solvent molecules
Polarons and bipolarons
FIGURE 16.3 Schematic representation of the reversible volume change associated with the electrochemical reactions of polypyrrole in electrolytes. (Modified from Otero, T.F., Polymer Sensors and Actuators, ed. D. de Rossi and Y. Osada, Springer-Verlag, Berlin, 2000. With permission.)
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H N
H N
H N
+
(a)
+ N H
H N
N H H N
H N
(b)
+ N H
N H
FIGURE 16.4 (a) Schematic representation of the chemical structure of neutral ideal lineal chain of polypyrrole, (b) schematic representation of the same partially oxidized chain (polaron or radical-anion). (Modified from Proceeding of SPIE, 397, 148, 2000.)
must be under control of the oxidation charge through the control of the counterion composition in the oxidized polymer (reaction 1). Reverse processes occur during the polymeric reduction by flow of a cathodic charge; the decrease of the film volume must be under the control of the cathodic charge through the control of the composition decrease of the progressively reduced film (reverse reaction 1). As the oxidized material is nonstoichiometric, the volume, as well as the material composition, changes continuously under the control of a continuous flow of charge (a constant direct current). The volume variation, like the electrochemical reaction, stops when the current is stopped and increases or decreases when the driving current increases or decreases. As we use these electrochemically generated volume-changes to produce a mechanical energy, the concomitant property is named electrochemomechanical. Structural shrinking and compaction of the film by cathodic polarization (anion prevailing interchange) [15–19], or by anodic polarization (cation prevailing interchange) [27], produce anomalous electrochemical responses in CPs [9,25] during swelling processes, which are simulated by the model [20–22]. Some groups of research try to explain the anomalous responses based on the cathodic compaction by a nonconducting–conducting transition during oxidation. These explanations are unable to describe anomalous responses based on anodic compaction processes observed during electrochemical reduction [27]. A different approach considers only the osmotic component to explain the swelling process [28]. These changes of volume have been detected in some of the pioneering works on the electrochemistry of conducting polymers [29–38] and later confirmed at the microscopic level by in situ AFM, ellipsometry, conductivity measurements, in situ electrogravimetry, and others [31,39–52]. Considering the existence of volume changes, Baughman et al. suggested the possibility for using conducting polymers as basic materials for the construction of actuators, similar to those developed from piezoelectric, electrostrictive, or ferroelectric inorganic actuators [53,54].
16.4
Basic Molecular Actuator and Muscle Similarity
The described theoretical model involves a basic molecular actuator constituted by an ideal and lineal polymeric chain. We can imagine this basic chain connected to a metallic electrode and immersed in an electrolyte (Figure 16.5). The strong intramolecular interactions produce a coil-compact structure of the chain. During oxidation, consecutive electrons are extracted from the chain, positive charges are
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(a)
Metal
5e−
Solution
ox
red
5e−
(b)
FIGURE 16.5 Molecular motor: Reverse conformational changes (mechanical energy) stimulated by oxidation or reduction of the polymeric chain in an electrolyte. (a) reduced chain, (b) oxidized chain. (From Otero, T.F., Modern Aspects of Electrochemistry, Kluwer Academic, New York, 1999. With permission.)
generated, and the above described processes now progressively transform the coil-like structure to a rod-like structure. The process can be stopped at any intermediate position, or reversed from any intermediate position, by stopping or reversing the direction of the current flow (the direction of the electrochemical reaction); this is the basic molecular motor working under electrochemical stimulation of the conformational movements and electrochemical control of the intramolecular interactions. Both the theoretical model and this basic molecular actuator include electric pulses, ions and water interchanges between the polymer and the solution, chemical reactions, stimulation of the conformational movements along polymeric chains, and changes in the inter- and intramolecular interactions. Those processes occurring in soft and wet materials mimic, at the molecular level, the consecutive events involved in the actuation of a natural anisotropic muscle.
16.5
Devices: Monolayers, Bilayers, Triple Layers, Combinations
Until now we have not been able to produce bundles of uniformly spaced single ideal chains of conducting polymers on electronic conducting material in order to reproduce a natural sarcomere (Figure 16.6) and perfect anisotropic lineal movements under electrochemical reactions. By chemical or electrochemical polymerization we produce films of branched, partially cross-linked (Figure 16.7) and, partially degraded materials (insoluble materials having different domains of charge storage capacity and conductivity). The molecular entanglement of these actual materials swells or shrinks under electrochemical oxidation or reduction, respectively. Both processes are three dimensional and quite isotropic. The way to translate these quasi-isotropic swelling and shrinking processes to a macroscopic movement was attained in 1992 by Swedish and Spanish teams independently, by the construction of
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Polymeric chains
Conducting material
FIGURE 16.6 Ideal structure of parallel lineal chains of conducting polymer attached to two nanometric lamellas of a conducting material in an electrolyte, mimicking a sarcomere. Under electrochemical oxidation the sarcomerelike structure swells and under reduction shrinks: Ideal anysotropic electrochemomechanical element.
FIGURE 16.7 Cross-linked structure of polypyrrole. (From Otero, T.F., Modern Aspects of Electrochemistry, Kluwer Academic, New York, 1999. With permission.)
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(a)
(b)
(c)
FIGURE 16.8 (a) Stainless steel electrode partially coated with an electrogenerated polypyrrole film, (b) adhesion of the nonconducting polymeric film on one side of the electrode, (c) the bilayer CP–tape is removed from the surface. (From Otero, T.F., and Sansin˜ena, J.M., Biolectrochem. Bioenerg., 38, 411, 1995.)
polymeric bilayers as indicated in Figure 16.8 [55–57]. A film of a conducting polymer was electrogenerated on a metallic electrode. Once rinsed and dried, a tape of a nonconducting polymer was attached to the conducting polymer film. The bilayer was peeled from the electrode and then the allorganic bilayer was used as a new electrode, by connecting the film of the CP to the working electrode (WE) of a potentiostat, in an electrolyte. The flow of reverse currents transforms the small macroscopic length variation (only a few micrometers) of the film, originated by the electrochemical reaction (Equation 16.1), into macroscopic angular (greater than 3608) movements (Figure 16.9). The rate of the electrochemical reaction follows the Butler–Volmer expression linking the reaction rate (the current density, i) and the electric potential, overpotential (h), that the CP is submitted to i ¼ i0 exp (aFh=RT )
(16:3)
FIGURE 16.9 Angular movement described by the free end of a bilayer muscle (CP–tape) under a current flow of 15 mA (a, b, and c), or of 15 mA (d and e), the muscle being immersed in a 0.1 M aqueous electrolyte.
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where a is the transfer coefficient (0a1), i0 is the exchange current density (i ¼ i0 ¼ kr cr [Cl]n exp(aFE0=RT), when h ¼ 0 and contains the reduction kinetic constant (kr), the concentration of the reduced polymer(cr) and the Cl concentration) for the studied material, and h is the applied overpotential (h ¼ EE0, E is the applied potential and E0 is the equilibrium potential) required to extract the electrons from the polymeric chains. The overpotential is a unique tool for the CP-based actuation; a shift of the overpotential for a few millivolts promotes a massive extraction of electrons from the material (anodic shift) or a massive injection of electrons (cathodic shift) with transformation of double bonds, conformational movements, interchange of ions and water molecules, etc. Other than these few millivolts, electromechanical actuators require tens to thousands of actuating volts. In addition, the flow of a constant current (i) produces an overpotential (h), which depends, through the electrochemical reaction (Equation 16.1), on the electrolyte concentration [Cl], the temperature of work (aFE0=RT), and the polymer composition (cr), providing unique properties of the electrochemomechanical actuators. In order to allow the current flow through the film of the CP, a metallic counter-electrode (CE) immersed in the electrolyte (Figure 16.10) is required. The use of a reference electrode (RE) allows perfect control of the CP potential (potentiostatic experiment), providing the potentiostat with the flow of the
CECE WE WE RE RE
I (mA)
∆E (mV)
e−
e− CP
Nonconducting polymer
Bilayer
Metal
Water discharge reaction
Solvent cations anions
Electrochemical reaction in CP
FIGURE 16.10 Electrochemical system checking a bilayer muscle (CP–tape). In order to allow the current flow through the muscle WE a metallic CE is required. So the charge flowing between the WE and the CE produces the electrochemical reaction generating the volume variation of the CP. The RE allows a perfect control of the muscle (CP) potential. (Modified from Otero, T.F., Polymer Sensors and Actuators, ed. D. de Rossi and Y. Osada, SpringerVerlag, Berlin, 2000. With permission.)
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required charge between the WE and the CE, to produce the electrochemical reaction (Equation 16.1) in the suitable extension and direction to attain the ordered potential according to the Nernst expression: E ¼ E0 þ RT =nF ln [(PPy nþ )s (Cl )n (H2 O)m ]=[PPy][Cl ]n
(16:4)
If [Cl]n, which represents the Cl concentration in solution, remains constant during the process, then the potential under stationary conditions depends only on the concentration relationships [(PPynþ)s(Cl) n (H2O)m]=(pPy) in the film. Experimentally we can submit the device to a potential step from E1 to E; this is a potentiostatic experiment. The potentiostat responds by producing the flow of a current (chronoamperogram is the recorded evolution of this current during and after the potential step). The [(PPynþ)s (Cl)n (H2O)m]=(PPy) relationship changes from that required when the potential was E1 to that required by the new potential E depending on the material. As any electrochemical transformation between (PPy) and [(PPynþ)s(Cl)n (H2O)m] generalizing from CP to oxidized CP promotes a volume variation of the film, a macroscopic angular movement of the free end of the bilayer artificial muscle is generated. The second most common way to control the movement of the device is by imposing from the potentiostat or galvanostat the flow of a constant direct current between WE and CE: a galvanostatic experiment. In this case, we can follow the evolution of the ‘‘muscle potential’’ (potential of the CP layer versus the RE) with time, recording the cronopotentiometric response. Whatever the methodology used to control the bilayer movement, a current flows through the conducting polymer of the bilayer in order to oxidize or reduce the polymer. The same current must flow through the metallic CE–electrolyte interface; the current can only flow if electrochemical reactions take place at the interface. Working in aqueous solutions, the most common of those reactions at a neutral pH are 2H2 O þ 2e $ H2 þ 2OH (when the CE acts as a cathode)
(16:5)
2H2 O $ O2 þ 4Hþ þ 4e (when the CE acts as an anode)
(16:6)
These reactions require a high overpotential. This means that an important fraction of the electrical energy applied to the device is consumed (overpotential times the current) by these reactions. An important fraction of the applied energy is wasted for producing changes in the solution pH (see OH [Equation 16.3] or Hþ [Equation 16.4] generation), inducing a chemical degradation of the conducting polymer. That means a reduction of the device lifetime. We note that the main components of the consumed electrical energy are electrochemical reactions in the PPy film, electrochemical reactions on the metallic CE, and Joule effects (ionic diffusion inside the film, ionic migrations, stimulated conformational movements of the polymeric chains, changes in the inter- and intramolecular interaction, and electrical connections—mainly at the metal wire–polypyrrole interface). In this context we can improve the efficiency of the consumed electrical energy, preventing at the same time the chemical transformation and degradation of the actuating material, by including the counter-electrode reaction as a useful part of the actuating device. This important problem was solved some time ago, as will be described later, by producing triple layers including two films of CP, each connected to one of the potentiostat–galvanostat outputs: the WE and the CE (Figure 16.11b) [55]. At present the construction of bilayers such as plastic–CP, paper–CP, or a thin film of any flexible material metal coated (i.e., by sputtering)–CP is very attractive [58,59] for most of the research groups involved in this field. The second problem of the laminar devices is that, in order to transform a quasi-isotropic change of volume into a macroscopic movement, we have mainly used the change in the PPy film length, ignoring longitudinal variations along the other two dimensions. This fact will again reduce the attained efficiency of the consumed electrical energy into mechanical energy. Nevertheless any solution to this problem will require perfect control of the synthesis of the above-proposed bundles or brushes
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a
b
c
d
e
WE
CE/RE
(a)
WE CE RE
WE
CE RE
Reference electrode Ag/AgCl
cathode
anode
nonconducting
film
anode
cathode
nonconducting
p Py films
film
Li+ClO−4
(b)
p Py films
Li+ClO4−
(c)
FIGURE 16.11 (a) Sequence followed for the construction of a triple layer muscle, (b) electrochemical equipment, electrical contacts, and scheme of the triple layer muscle (CP–two sided tape–CP). As the current flows, the CP film acting as anode swells and the second film of CP acting as cathode shrinks. The free end of the muscle (the other end supports the adhered electric contacts) describes the angular movement indicated by arrows. By reversing the direction of the current, the movement occurs in the opposite direction. The overall muscle potential (WE versus CE) is followed during the current flow. The muscle works in aqueous solution. (c) Distribution required for potentiostatic control, or during galvanostatic control, of the muscle in order to follow the potential of the CP acting as WE. (From Otero T.F., and Corte´s M.T., Adv. Mat., 15, 279, 2003. With permission.)
of ideal linear molecular-based motors. Their assembling to mimic sarcomeric units will allow the construction of basic perfectly anisotropic and highly efficient electrochemomechanical motors. Until then different research lines are open; the synthesis of the materials can be improved, the theoretical models can be developed, and different devices suitable for medical or technological applications can be designed. Different research groups are focused on the production of asymmetric monolayers of conducting polymers by physical means such as growing the CP on adsorbed and porous materials; or by chemical means such as generating a film of CP with a concentration gradient of counterions or of cross-linking network; or by generating a bilayer of the conducting polymer with a macroanion (shrinks by oxidation), and then generating a second layer of the same conducting polymer with a small anion (swells by oxidation), or even placing a metal sheet between both films [60–73]. These asymmetric films also
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produce a bending movement and need an electrolyte and a counterelectrode in order to allow the flow of the current and to produce the bending movement. This counter electrode will generate similar problems to those using bilayers as described earlier.
16.6
Linear Displacement Actuators
Many groups expend a lot of effort in order to manufacture devices directly producing linear displacements. Fibers and films of conducting polymers were directly used to follow their electrochemomechanical properties as will be pointed [74–101]. A second approach consists on the electropolymerization of a conducting polymer on helical metallic wires to generate a tube, or on zigzag metal wires to generate films, etc. [102–107]. Those efforts are based on a general belief on a conducting–nonconducting transition during the reduction of the conducting polymers, which should originate a drop in the electrical potential along the polymer, and a nonuniform actuation. The supporting metal wire should guarantee a uniform potential and current distribution. But the muscles always work under partial oxidation of the material. These actuators also need a metallic counter electrode to allow the current flow. Moreover, for such thick films as those usually employed for the construction of macroscopic actuators the experimental results indicate, and the ESCR model predicts, a very difficult and time consuming reduction completion at very high cathodic overpotentials [10,11]. The polymeric entanglement is closed under partial reduction, entrapping counterions and keeping a high enough conductivity (>102 S cm1) to allow a uniform actuation not requiring any embedded metallic wire for conducting reasons. Digital image analysis of the bending movement also indicates uniform bending of the laminar full polymeric devices irrespective to the distance to the metallic electrical contact [108,109]. Both monolayers and bilayers include a new paradigm related to traditional inorganic actuators (piezoelectric, electrostrictive, ferroelectric, electrokinetics, etc.); the electrode material is, simultaneously, the actuator and the support for the current flow, which is the driving force for the electrochemical reaction producing the volume variation of the material. Moreover, actuation requires the presence of an electrolyte as the counterion supplier. All the inorganic actuators (and the new polymeric ones mimicking them) use two inert metallic electrodes in order to produce and to apply an electric field on the actuating material. That means that all the electromechanical devices are constituted by three layers: metal electrode–actuating polymeric material–metal electrode. It is not possible to imagine any inorganic actuator, constituted by an isolated monolayer of a piezoelectric or electrostrictive material, to also act as the electrode to support the electric field. The electrochemical synthesis of the films of conducting polymers, and the electrochemical actuation, are suitable for the construction of elegant and imaginative microdevices and microtools bilayers by using microelectronic technologies [110–121].
16.7
Triple Layer
Using a two-side tape to construct a CP–tape bilayer, the free side of the tape is adhered to a second film of the same CP, previously electrogenerated on a metallic electrode (Figure 16.11a). The triple layer CP–tape– CP is then removed from the metal and connected to the potentiostat as indicated in Figure 16.11b and Figure 16.11c. Any of the methodologies include WE and CE in the same device. When one of the CP films acts as an anode (Equation 16.1 forward), which swells and pushes the device, the second CP film acts as a cathode (Equation 16.1 backward), which shrinks and trails the device. The device includes the actuating electrochemical reactions from both electrodes (WE and CE); when one of the films swells producing a pushing force, the second film shrinks generating a trailing force. No metallic counterelectrode is required. Both electrochemomechanical reactions, volume generation (at the anode or at the CP working electrode) and volume destruction (at the cathode or at the CP counter electrode) are actuating reactions, improving the efficiency of the consumed electrical energy and avoiding both pH variations of the media and chemical degradation of the actuating materials. Moreover, the different oxidation states of both actuating
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electrodes define two stationary potentials (Ei and Ej) generating a battery (DE ¼ Ei Ej). When the work of the muscle promotes an increment of the muscle potential, DE increases and we must provide an external electrical energy. But when the actuation promotes a decrease in DE we can recover from the muscle a maximum electrical energy (improving again the efficiency) of DG ¼ nFErev, where n is the number of electrons transferred per polymeric segment, F is the Faraday constant, and Erev is the potential of both electrodes when the battery is discharged. The recovered self-discharge current is given by i ¼ i0 [ exp (a1 z1 FDE=RT ) exp ( (1 a2 )z2 FDE=RT)
(16:7)
where i0 is the interchange current of any of the polymeric electrodes when DE is zero, a1 and a2 are the transfer coefficients for the oxidation and reduction process of the nonstoichiometric compounds. The device (also now the battery) potential DE decreases with time. The methodology from Figure 16.11b allows the control of the electrochemical reaction rates and electrochemomechanical property variations (swelling and shrinking processes linked to the forward and backward reactions) taking place in the device through the applied current. Simultaneously we can follow the muscle potential (WE versus CE) using only two connecting wires. This fact will later become important to describe intelligent (conscious) devices. The methodology described by Figure 16.11c allows a chronoamperometric control of the device by applying a constant potential to the CP film acting as WE and recording the evolution of the current flowing through the system to attain the desired potential and the concomitant new position of the device. Working under constant current we can follow the evolution of one of the PPy film potential acting as WE. Considering the overall muscle potential measured using the methodology represented in Figure 16.11b, we can separate both half potential components (for each of the constituent CP films). The structure of the triple layer is suitable for the construction of muscles working in air, by encapsulation of the triple layer keeping a small amount of electrolyte, or by using an adherent, flexible, and ionic conducting membrane between both films of the conducting polymers [55,91,122–127].
16.8
Characterization of the Electrochemomechanical Muscles (Working Potential Ranges)
One of the first points we need to know in order to characterize any actuator based on conducting polymers is the range of electric potentials where the swelling or shrinking (oxidation or reduction) processes occur. A PPy–tape bilayer, using the experimental methodology from Figure 16.11c, is submitted to consecutive cyclic potentials between two potential limits until the obtension of a stationary voltammogram. The potential limits and the rate of the potential sweep are changed to obtain the overall oxidation and reduction of the material (Figure 16.12), the maximum charge under the voltammogram. At very low sweep rates (<1 mV=s) the oxidation processes for polypyrrole muscles in LiClO4 aqueous solution occur from 950 mV, versus SCE, to 500 mV. At cathodic potentials higher than 950mV, the closed polymeric structure (partially oxidized) is compacted by slow reduction under conformational relaxation control. At higher anodic potentials than 500 mV, the material oxidation coexists with water discharge and polymeric degradation. So under very slow sweep rates, the optimum oxidation or reduction and swelling or shrinking processes occur between these two potentials, consuming an oxidation charge Qox that depends on the film weight. When the potential sweep rate is increased, both oxidation and reduction maxima shift toward more anodic or cathodic potentials, respectively. These shifts are linked to the slower of the processes involved in the electrochemical reaction (rate determining step). The right identification of the slower process (conformational changes, diffusion processes, migration processes, ohmic resistances, etc.) is a basic point for improving the device’s properties. Anyway, very high anodic potentials can be attained (up to 10 V) for short periods of time without any major degradation of the material; all of the flowing charge is consumed to complete the oxidation of the conducting polymer film [128,129]. On the other hand,
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I (mA/cm2)
Qox 0 Qred
E1
E2
E(mV)
FIGURE 16.12 Voltammogram obtained from a polypyrrole film in 0.1 M LiClO4 aqueous solution at 20 mV s1, using a platinum sheet as counter electrode. Oxidation and swelling processes start on E1, when the anodic current rises for increasing anodic potentials. Beyond the potential E2, degradation polymeric processes occur. These processes are inhibited meanwhile, at high potential sweeps or by flow of high currents, the oxidation charge (Qox) is not completed, even if high potentials such as 10 V is attained. The polymer shrinks during reduction from E2 to E1 and compacts when, at very low reduction rates E1 is attained
very high cathodic potentials can be attained for short periods of time without any polymeric compaction; the film still remains very oxidized and swelled, all the cathodic charge being consumed to reduce and shrink the material. So when we construct muscles using these materials, the muscles can be checked potentiostatically, by a potential step to a constant potential lower than 500 mV for very long periods of time, or galvanostatically (by flow of a constant current through the electrode) for long periods of time if the film acting as anode does not exceed 500 mV. If it does, the material is degraded by water discharge and hydroxylation of the polarons. Nevertheless, if high potentials are attained for short periods of time and meanwhile the material is partially oxidized, no degradation is observed. In the case of the muscles, these good conditions are maintained, whereas the muscle movement is far away from the limits expected for its angular movement. If the rate-limiting step for the electrochemical reaction responsible for the slow movement is attributed to the ion diffusion in solution, the use of ionic liquids during synthesis and for the device control must facilitate the movement at low muscle potentials [85,126,130–139]. Longer lifetimes were obtained without any significant rate increase or muscle potential decrease.
16.9
Sensing and Actuating Capabilities of Those Electrochemical Reactions
Equation 16.1 can be rewritten considering these main variables acting on each of the reaction component: k
(PPy)s þ n(Cl )aq þ mH2 O $[(PPy nþ ) (Cl )n (H2 O)m ]gel þ (ne )metal T
(Volume)1 (conc: and I) (synthesis
s
(Volume)2
i and E
conditions) This indicates the existence of a volume change (V1$V2) during the oxidation or reduction processes that can be influenced by the concentration of counterions in solution, whose ionic strength I (as well as that of the free water activity) is influenced by the presence of other ions (acting or not on the
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electrochemical reaction). The constant rates of the forward and backward reactions (k) are influenced, through the Arrhenius expression, by changes on the experimental temperature. In addition, V1 and V2 depend on the conditions used for the polymer synthesis, mainly on those controlling the cross-linking and degradation processes. When different experiments are performed at a constant current under different values of any variable acting in the reaction (the above mentioned or different external pressure or mechanical stress: they act on the volume change), the evolution of the muscle potential (E) during the movement must be influenced by the variable. This contains a unique possibility: the actuating device under the control of the current can sense, through E, any variable (such as those indicated above) acting on the electrochemical reaction.
16.10
Actuating Capabilities of an Artificial Muscle
0.25 0.2 0.15 0.1 0.05 0
14
21 28 35 Current (mA)
Electric charge (mC)
(a)
7
(c)
600 500 400 300 200 100 0
Electric charge (mC)
Movement rate (rad/s)
Volume changes between (PPy)s and [(PPynþ)s(Cl)n (H2O)m]gel are under the main control of the amount of Cl anions (amount of counterions, for any electrolyte) interchanged with the solution. Figure 16.13a shows the expected linear relationship between the current and the rate of the angular movement; the movement is under the control of the current. Moreover, according to the electrochemical reaction, the consumed charge controls the variation of the counterion concentration, the volume change, and the angle (Figure 16.13c) described by the artificial muscle [12,140–143]. Moreover, whatever the testing current, a constant charge (Figure 16.13b) was consumed to describe the same angle. Figure 16.13 confirms the electrochemical nature of the movement and the stated relationship between position and charge. We can define the charge required to describe an angular movement of
360 300 240 180 120
7
14
(b)
21 28 35 Current (mA)
8 10 15 20 25 30 35
0
50 100 150 Angle (degrees)
200
FIGURE 16.13 (a) Linear relationship between the applied current and the angular rate determined from the times required to describe an angular movement of 908 by a triple layer muscle (2 cm 1.5 cm 13 mm), constituted by two polypyrrole films weighing 6 mg each, and checked under different current densities (8, 10, 15, 20, 25, 30, and 35 mA) in a 1 M LiClO4 aqueous solution, (b) electric charges consumed through 908 under different current densities. (c) Electric charge consumed by the triple layer muscle to describe different angles (308, 458, 608, 908, 1208, 1358, and 1808) under the currents studied. Experiments in 1 M LiClO4 aqueous solution. (From Otero T.F., and Corte´s M.T., Chem. Commun., 3, 284, 2004. With permission.)
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8
Current (mA/mg)
7 6 5 4 3 2 1 0 0
0.1
0.2
0.3
0.4
0.5
0.6
Angular rate (rad/s)
FIGURE 16.14 Angular rate measured through a movement of 908 using triple-layer muscles of different dimensions (including different weights of polypyrrole films: 8.3, 7.8, 7.4, 6, 5.5, 5.1, 5, 4, 3.7, 3.5, 3, 2.3, and 2 mg) and checked in 1M LiClO4 aqueous solution under different currents (10, 15, 20, 25, and 30 mA). (From Otero T.F. and Corte´s M.T., Chem. Commun., 3, 284, 2004. With permission.)
18 (a mC=degree), and using it we can obtain how much charge, and the direction of flow, is required to attain any new position (any new angle) from the actual position: Q(mC) ¼ a(mC=degree) angle(degrees)
(16:8)
This is an electrochemopositioning device, requiring a constant charge to describe a defined angle, i.e., 908; the time consumed for the movement must be under the control of the charge flowing through the device per unit of time, i.e., the current [128,142]. This electrochemical nature is underlined in Figure 16.14; different artificial muscles, having different surface area or constructed with PPy films having different thicknesses (different weights of polypyrrole), produce the same angular-movement rate under flow of analogous charge per unit of time (current) and per unit of CP weight (same variation of the oxidation deep). So the electrochemical nature of the movement allows a perfect control of both the movement rate (by the current) and the angular position (by the charge). This means we have a perfect machine able to transform electrical energy into mechanical energy using as mechanical components the electrochemically stimulated conformational movements of the polymeric chains. The currently available electrical generators are able to give very precise electrical currents and soft current variations, so we can produce with our artificial muscles very precise movement rates and soft rate variations, mimicking elegant movements such as those from mammals.
16.11
Sensing Capabilities of an Artificial Muscle
As expected from the electrochemical reaction (Equation 16.1), the muscle potential is influenced by the chemical and physical variables [144–150]. Figure 16.15a through Figure 16.15d shows the evolution of the muscle potential (CP of the WE versus CP of the CE) from a triple layer under different concentrations of the electrolyte, different temperatures, different weights attached to the bottom of the muscle, or different current flowing through the device. Figure 16.15a through Figure 16.15d shows the linear evolution of the electrical energy consumed by the artificial muscle as a
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1.6
log(Ee) (kJ/kg)
1.4
1.2
1
0.8
0.6 −1.5
−1
−0.5
0
0.5
Log [LiClO4] (M)
(a)
15
Ee (kJ/kg)
14 13 12 11 10 9 0
10
20
(b)
30
40
50
60
T (⬚C) 300 Y = 59 + 6.5 X r = 0.99
Ee (mJ)
250 200 150 100 5 (c)
10
15
20 25 i (mA)
30
35
40
FIGURE 16.15 A triple layer (2 1.5 cm2, 12 mg of PPy) describes 908 in aqueous solutions of LiClO4 (3, 1, 0.5, 0.25, 0.1, and 0.05 M) under a constant current of 10 mA, in 0.1 M LiClO4 at different temperatures: 58C, 158C, 258C, 358C, and 458C, Under flow of different currents: 5, 10, 15, 20, 25, and 30 mA. (continued)
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80
Ee (mJ/mg)
70 60 50 40 30 20 10 0
500
(d)
1000 1500 Load (mg)
2000
2500
FIGURE 16.15 (continued) Consumed electrical energy by the device as a function of the different studied variables: (a) electrolyte concentration, (b) temperature, (c) current, and (d) shifted weight.
function of the studied experimental variable, one by experimental series, and as was expected from the relationship between variables involved in Equation 16.1. Rising driving currents or higher shifted weights produce higher overpotentials consuming increasing electrical energy. Rising electrolyte concentrations or temperatures require lower overpotentials consuming lower electrical energy during the movement. These results underline the simultaneous sensing and actuating capabilities of the device. Both signals, the actuating current and the muscle potential response, are included in the same two connecting wires (Figure 16.11b) opening a new paradigm for the robotic devices.
16.12
Tactile Sensibility
The ability of the devices for sensing mechanical variables opens the way for exploring one of the most exciting possibilities that scientists are trying to reproduce—the tactile sense. Considering that trailing increasing weights by the muscle we obtain increasing muscle potentials, if a muscle moves freely and meets an obstacle opposing a mechanical resistance lower than the mechanical energy produced by the device, the device must touch, push, and shift the obstacle (Figure 16.16). At the beginning of the movement, the muscle moves freely until it touches the obstacle; the evolution of the muscle potential must be the same as that of the muscle moving without any hanging weight. When the muscle touches the obstacle it feels a mechanical resistance and, according to the above conclusion, the muscle produces an extra energy by raising its overpotential [8,151,152]. Depending on how fast this evolution occurs, we can obtain a response to the touching moment. Figure 16.17 shows that there exists an instantaneous evolution of the potential (the response is a potential step). Moreover, if we repeat the experiment using obstacles producing increasing resistances (increasing weight obstacles), increasing potential steps are produced, proportional to the obstacle weight at the moment of the contact; this is a touching sensor. The moving device indicates the moment of contact and the mechanical resistance opposed by the obstacle. When the mechanical resistance of the obstacle exceeds the mechanical energy produced by the device, the device is unable to shift the obstacle and the muscle potential steps to very high values at the moment of contact. So we have a muscle with tactile sense: quite simple response-analysis software can transform the ensemble (computer, potentiostat, and device) into a conscious system. The system indicates when a muscle, or the mechanical tool driven by the muscle, touches an obstacle and how much mechanical resistance the obstacle opposes.
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(a)
(d)
(b)
(e)
(c)
(f)
FIGURE 16.16 (a) The triple layer muscle initiates its movement under a constant current of 5 mA, in 1 M LiClO4 aqueous solution; 10 s later (b) the muscle meets the obstacle weighing 6000 mg, pushing and sliding it (c and d). (e) The angular movement allows the muscle to overcome the border of the obstacle. (f) The free movement goes on until the current stops.(From Otero T.F., Adv. Mat. 15, 279, 2003. With permission.)
j:14400 mg i:9600 mg h:8400 mg g:7200 mg f:6000 mg e:4800 mg d:3600 mg c:2400 mg b:1200 mg a:non obstacle
3
Potential (V)
2.5 2 1.5 1 0.5 0 −0.5 0
10
20
30
40
50
60
70
Time (s)
FIGURE 16.17 Chronopotentiograms obtained from a triple layer macroscopic muscle containing two polypyrrole films [2 cm 1.5 cm 13 mm] weighing 6 mg each one, under flow of 5 mA in 1 M LiClO4. The muscle moves freely, contacting an obstacle after 10 s and slides it for 3.5 s, overcomes its border and continues with a full angular movement of 1088 (from 188 to þ908). The initial position is recovered by applying a current of 5 mA for 57 s. Obstacles weighing 1,200, 2,400, 3,600, 4,800, 6,000, 7,200, 8,400, 9,600 mg were slid, but the muscle was unable to push and slide an obstacle weighing 14,400 mg. (From Otero T.F., Adv. Mat. 15, 279, 2003. With permission.)
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Conjugated Polymers: Processing and Applications
Mechanical Characterization of Electrochemomechanical Muscles
Electrochemically induced volume variations in CP produce length variations and forces. Both variables can be quantified by using standard tensile testing machine or specific lever or cantilever machines constructed for this aim. The samples are individual films of CP or devices (bilayer, triple layers, or monolayers; free-stand samples; or applied on metallic, helical, or zigzag wires). The sample must be totally or partially immersed in an electrolyte. Measurements can be performed by keeping a constant length of the sample and following the force variation induced by the electrochemical reactions, or by applying a constant force to the sample and following the length variation induced by the electrochemical reactions [74–107]. The force–displacement data are transformed to strain–stress curves, which allow the estimation of the Young’s modulus. In this way properties of the system such as stress, strain, strain rate, work density, power to mass ratio, coupling, efficiency, and life time (cycle life) can be quantified (Table 16.1) and compared from different materials and systems (Table 16.2). The influence of different variables: electrolyte concentration, electrolyte pH, electrochemical driving method, etc., on the properties of the different samples and devices can be followed. Moreover, results claiming the development of new lineal devices developing giant strains of 12%, 20%, and up to a 26% of the original length of the actuator, but dropping very fast to much lower strains with electrochemical cycling, are being described [84,88,93]. These results show the great technological possibilities of CP for the development of nanoscopic, microscopic, and macroscopic actuators. A lot of expectancies emerge for engineers and physicists. The mechanical characterization of materials and devices, altogether the construction of microdevices, and the exploration of new materials, synthesized in different electrolytes or using ionic liquids, are attracting the interest of most of the researchers involved in this fascinating field. In contrast, the popularity of these polymeric materials from the point of view of a mixed theoretical treatment by the polymer science and the electrochemistry and their electrochemomechanical characterization is very low. Even if most of the interest is focused on the development of new materials and on new ways for improving their characteristics as new electrical motors, some stationary state had been attained several years ago, hindering the construction of artificial muscles of any shape and volume, and having any mechanical energy . This apparently classical problem has hindered for more than a decade the development of macroscopic muscles, sensors, and actuators based on CP.
TABLE 16.1
Mechanical Characteristics of the Conducting Polymers as Electrochemomechanical Actuators
Property Strain (%) Stress (MPa) Work density (kJ=m3) Strain rate Power (W=Kg) Life (cycles) Coupling Efficiency (%) Modulus (Gpa) Tensile strength (Mpa) Applied potential (V) Charge transfer (C=m3) Conductivity (S=m) Cost ($=kg)
Minimum
Type 2 5 100 1 28,000
0.2
<1 0.8 30 1.2
107 10,000 3
Source: From Madden, J.D.W., IEEE J Oceanic Eng., 29, 706, 2004. With permission.
Maximum 12 34 12 150 800,000 0.1 18 3 120 10 108 45,000 1,000
Limit >20 200 1,000 10,000 1,00,000
400
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TABLE 16.2 Comparison between Two Mechanical Properties of Different Actuating Materials: Skeletal Muscles, Thermomechanical (Thermal Liquid Crystals and Thermal Shape Memory Alloys), Electrochemomechanical (Conducting Polymers and Carbon Nanotubes) and Electromechanical (Ionic Polymer Metal Composites, Field Driven Liquid Crystal Elastomers, Dielectric Elastomers) Material
Work Density (kJ=m3)
Strain (%)
<40 5.5 10a 20a 40 56 100b 100 150a 320a 1000
20 0.5 120 max 2 0.2 19–45 6 2 380 max 3.5 5
Skeletal muscle IPMC Dielectric elastomer (silicone) Field driven liquid crystal elastomer[31] Carbon nanotubes Thermal liquid crystal Ferromagnetic SMA Conducting polymer Dielectric elastomer(VHB) Ferroelectric polymers Thermal SMA
Source: From Madden, J.D.W., IEEE J Oceanic Eng., 29, 706, 2004. With permission. a Represents internal energy density that must be coupled to the load. Coupling will reduce energy density by a factor of 2 at least. b Does not include the coil volume.
16.14
New Expectancies
Besides, the electrochemical stimulation of the conformational movements and considering that the term artificial muscles is currently accepted for any transduction between any stimulating energy and a mechanical work through conformational changes in polymeric molecules, different artificial muscles can be envisaged based on conducting polymers. Taking into account the presence of dipole moments in the polymer chain and the photochemical or magnetic excitation of the solitons present in the chain, electromechanical, photomechanical, or magnetomechancial actuators can be envisaged. So, theoretical calculations predict the existence of photostimulation [153]. The soliton-induced conformational effects create sufficient strains to provide an intrinsic high-strain-rate actuation mechanism in optical excitation processes. The advantage of this mechanism should be that slow counterion processes are unnecessary for obtaining very fast conformational movements. If those theoretical calculations can be experimentally checked, a new field of ultra-fast photomechanical devices will be soon explored.
16.15
Products and Companies Based on CP Artificial Muscles
The described state of the art can be historically followed through some overview references [12– 14,53,109,128,141,153–160], and drives the coming exploration through two different ways: (a) development of a multielement using basic lineal actuators and a metallic counter electrode (Figure 16.18a and Figure 16.18b) and (b) development of basic elements describing both bending and linear movements, each basic element constituted by triple layers including two (Figure 16.19a and Figure 16.19b) actuating electrodes (electrode and counter electrode) able to simultaneously record the muscle potential (the CP acting as CE is short-circuited to the RE) under the influence of any experimental variable. Those basic elements (microscopic or macroscopic) can be connected in order to produce a basic three-dimensional unit (Figure 16.20) keeping the electrochemomechanical characteristics of each constituent basic element. Now, by repetition (n times) of the basic unit we could construct an artificial muscle of any shape and any volume and supply any amount of
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(a)
1 piece
10 pieces (b)
FIGURE 16.18 (a) Schematic diagram of some bundles of PPy–metal coil composites. (b) Photographs of a PPy–W coil composite actuator and a bundle of its 10-piece actuator. (From Hara, S., Zama, T., Takashima, W., Kaneto, K., Synth. Met., 146, 47, 2004. With permission.)
mechanical energy (n times the energy produced by the basic unit). Up to now one of the main problems to construct the basic unit is not focused on the suitability of the CP, or on the efficiency to transform electrical energy into mechanical energy, or on the electrochemomechanical properties of the materials, as can be expected. The limiting step is very basic and very primitive: the nonuniformity of the electrical connections between CP and metallic wires promotes irregular ohmic drops and irregular current distributions across the unit giving chaotic relative-movements of the different basic elements. A similar problem, now related to the irregular distribution of the electric field and the mechanical resistance of the system, is hindering the development of robotic tools based on electromechanical devices.
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(a)
(CE)
(WE)
(1)
Oxidation of WE Reduction of CE
(3)
Reduction of CE Oxidation of WE
(4) (5) (6) (2)
(b)
FIGURE 16.19 (a) Basic element constituted by two triple layers moving by opposition: during the observed contraction the two internal layers of CP act as cathode (CE short-circuited to the RE) and the two external ones as anode. The original position is recovered by changing the sense of the current flow. All the time we follow the potential of the overall element. (b) Basic element constituted by four bilayers including WE, CE, and RE suitable to follow the device potential at any moment.
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FIGURE 16.20 Basic three-dimensional electrochemomechanical unit constituted by 36 basic element, i.e., those from Figure 16.19a or Figure 16.19b, microscopic or macroscopic elements. By repetition of this basic unit a muscle of any shape, volume, and supplying any mechanical energy can be constructed.
If we look at the Web, a lot of companies claim the development of basic elements for the manipulation of medical tools at the end of probes or for driving those catheters or endoscopes along arteries or along pipes containing different fluids, the construction of artificial sphincters, Braille screens, etc., using artificial muscles. In most of the cases a clarification about the nature of these named artificial muscles does not exist. Nevertheless, these facts, together with the existence of a priority economical support from DARPA (USA) and ESA (EU) agencies of the polymeric actuator area, have established a strong competition to get funds, to develop new devices, and to get products. Professor Kaneto in Japan has started a company EAMEX, mainly devoted to the optimization of CPbased actuators. The Australian Intelligent Polymer Research Institute (Prof Wallace) is collaborating with different companies like Quantum Technology Pty. Ltd (Development of Braille screens) or with the Cooperative Research Center for Cochlear Implant and Hearing Aid Innovation. Micromuscles AB in Sweden and Molecular Mechanisms in Massachusetts, always working in collaboration with academic groups (Professors Inganas, Smela, Madden, DeRossi, etc.) are developing different medical tools or devices for medical applications. Many more companies can be found in the Web claming by the development of artificial muscles based on conducting polymers, but going inside the page it is very difficult to differentiate among wishes and realities. Anyway, after 13 years of academic research the time for the first important stage of
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technological developments has arrived. A great success is expected for the development of threedimensional simultaneous sensor–actuator devices if we are able to collect synergies from the electrochemical, polymeric, and mechanical aspects of those fascinating devices. A new world of soft and wet, sensing and tactile (conscious) actuating machines is expecting for the required intellectual energy and engineering ability of the young scientists.
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46. Bartlett, P.N., P.R. Birkin, M.A. Ghanem, and C.S. Toh. 2001. Electrochemical syntheses of highly ordered macroporous conducting polymers grown around self-assembled colloidal templates. J Mater Chem 11:849. 47. Juttner, K., and C. Ehrenbeck. 1998. Electrochemical measurements of the ion conductivity, permselectivity and transference numbers of polypyrrole and polypyrrole derivatives. J Solid State Electrochem 2:60. 48. Chen, X.W., K.Z. Xing, and O. Inganas. 1996. Electrochemically induced volume changes in poly(3,4-ethylenedioxythiophene). Chem Mater 8:2439. 49. Ehrenbeck, C., and K. Juttner. 1996. Ion conductivity and permselectivity measurements of polypyrrole membranes at variable states of oxidation. Electrochim Acta 41:1815. 50. Schmidt, V.M., D. Tegtmeyer, and J. Heitbaum. 1995. Transport of protons and water through polyaniline membranes studied with online mass-spectrometry. J Electroanal Chem 385:149. 51. Ryan, M., E. Bowden, and J. Chambers. 1994. Dynamic electrochemistry—methodology and application. Anal Chem 66:R360. 52. Yang, H., and J. Kwak. 1997. Mass transport investigated with the electrochemical and electrogravimetric impedance techniques. 1. Water transport in PPy=Cupts films. J Phys Chem B 101:774. 53. Baughman, R.H., L.W. Shacklette, R.L. Elsembaumer, E. Plitcha, and C. Becht. 1990. Conducting polymer electromechanical actuator. In Conjugated polymer materials. Opportunities in electronics, optoelectronics and molecular electronics, eds. J.L. Bre´das and R.R. Chance. Netherlands: Kluwer Academic Publisher. 54. Baughman, R.H., and L.W. Shacklette. 1991. Science and Application of Conducting Polymers, ed. W.R. Salaneck, D.T. Clark, and E.J. Samuelson, 47. Bristol: Adam Hilger. 55. Otero, T.F., et al. 1992. Patents: ES 2 048 086 and ES 2 062 930. 56. Otero, T.F., E. Angulo, J. Rodrı´guez, and C. Santamarı´a. 1992. Electrochemomechanical properties from a bilayer: Polypyrrole=non-conducting and flexible material. Artificial muscle. J Electroanal Chem 341:369. 57. Pei, Q.B., and O. Inganas. 1992. Conjugated polymers and the bending cantilever method— electrical muscles and smart devices. Adv Mater 4:277. 58. Deshpande, S.D., J. Kim, and S.R. Yun. 2005. New electro-active paper actuator using conducting polypyrrole: Actuation behaviour in LiClO4 acetonitrile solution. Synth Met 149:53. 59. Higgins, S.J., K.V. Lovell, R.M.G. Rajapakse, N.M. Walsby. 2003. Grafting and electrochemical characterisation of poly-(3,4-ethylenedioxythiophene) films, on Nafion and on radiation-grafted polystyrenesulfonate–polyvinylidene fluoride composite surfaces. J Mater Chem 13:2485. 60. Wang, H.L., J.B. Gao, J.M. Sansin˜ena, and P. McCarthy. 2002. Fabrication and characterization of polyaniline monolithic actuators based on a novel configuration: Integrally skinned asymmetric membrane. Chem Mater 14:2546. 61. Sansinena, J.M., J.B. Gao, and H.L. Wang. 2003. High-performance, monolithic polyaniline electrochemical actuators. Adv Funct Mater 13:703. 62. Onoda, M., H. Shonaka, and K. Tada. 2005. A self-organized bending-beam electrochemical actuator. Curr Appl Phys 5:194. 63. Okamoto, T., K. Tada, and M. Onoda. Bending machine using anisotropic polypyrrole films. Jpn J Appl Phys 1 39(2000): 2854; Jpn J Appl Phys 2 38(1999): L1070. 64. Shakuda, S., S. Morita, T. Kawai, and K. Yoshino. 1993. Dynamic characteristics of bimorph with conducting polymer gel. Jpn J Appl Phys 1 32:5143. 65. Onoda, M., Y. Kato, H. Shonaka, and K. Tada. 2004. Artificial muscle using conducting polymers. Electron Eng Jpn 149:7. 66. Onoda, M., and K. Tada. 2004. Anisotropic bending machine using conducting polypyrrole. IEICE T Electron E87C: 128. 67. Okamoto, T., Y. Kato, K. Tada, and M. Onoda. 2001. Actuator based on doping=undoping-induced volume change in anisotropic polypyrrole film. Thin Solid Films 393:383.
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