Polymer electrolytes
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Polymer electrolytes Fundamentals and applications Edited by César Sequeira and Diogo Santos
Oxford
Cambridge
Philadelphia
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© Woodhead Publishing Limited, 2010
Published by Woodhead Publishing Limited, Abington Hall, Granta Park, Great Abington, Cambridge CB21 6AH, UK www.woodheadpublishing.com Woodhead Publishing, 525 South 4th Street #241, Philadelphia, PA 19147, USA Woodhead Publishing India Private Limited, G-2, Vardaan House, 7/28 Ansari Road, Daryaganj, New Delhi – 110002, India www.woodheadpublishingindia.com First published 2010, Woodhead Publishing Limited © Woodhead Publishing Limited, 2010 The authors have asserted their moral rights. This book contains information obtained from authentic and highly regarded sources. Reprinted material is quoted with permission, and sources are indicated. Reasonable efforts have been made to publish reliable data and information, but the authors and the publisher cannot assume responsibility for the validity of all materials. Neither the authors nor the publisher, nor anyone else associated with this publication, shall be liable for any loss, damage or liability directly or indirectly caused or alleged to be caused by this book. Neither this book nor any part may be reproduced or transmitted in any form or by any means, electronic or mechanical, including photocopying, microfilming and recording, or by any information storage or retrieval system, without permission in writing from Woodhead Publishing Limited. The consent of Woodhead Publishing Limited does not extend to copying for general distribution, for promotion, for creating new works, or for resale. Specific permission must be obtained in writing from Woodhead Publishing Limited for such copying. Trademark notice: Product or corporate names may be trademarks or registered trademarks, and are used only for identification and explanation, without intent to infringe. British Library Cataloguing in Publication Data A catalogue record for this book is available from the British Library. ISBN 978-1-84569-772-3 (print) ISBN 978-1-84569-977-2 (online) The publisher’s policy is to use permanent paper from mills that operate a sustainable forestry policy, and which has been manufactured from pulp which is processed using acid-free and elemental chlorine-free practices. Furthermore, the publisher ensures that the text paper and cover board used have met acceptable environmental accreditation standards. Typeset by Toppan Best-set Premedia Limited, Hong Kong Printed by TJI Digital, Padstow, Cornwall, UK
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Contents
Contributor contact details Preface
xi xv
Part I Types and development of polymer electrolytes
1
1
3
1.1 1.2 1.3 1.4 1.5 1.6 1.7 2 2.1 2.2 2.3 2.4 2.5 2.6 2.7 2.8 2.9 2.10
Introduction to polymer electrolyte materials C. A. C. Sequeira and D. M. F. Santos, Technical University of Lisbon, Portugal Introduction Categories of polymer electrolytes Structure and its implications Conductivity measurements Applications in practical devices Conclusions References Ceramic polymer electrolytes J. S. Syzdek, Warsaw University of Technology, Poland Introduction Experimental approaches First composites – conductive fillers Development of insulating fillers Impact of the filler surface on the transport properties Interfacial concerns Other types of ceramic–polymer systems Conclusions Acknowledgements References
3 5 17 20 39 51 52 62 62 67 70 72 75 77 80 82 83 84 v
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3
Polymer electrolytes based on natural polymers A. Pawlicka and J. P. Donoso, Universidade de São Paulo, Brazil Introduction Grafted natural polymer-based solid polymer electrolytes Plasticized natural polymer-based solid polymer electrolytes Other natural polymer-based systems Magnetic resonance spectroscopy of polymer electrolytes obtained from natural polymers Conclusions and future trends References
3.1 3.2 3.3 3.4 3.5 3.6 3.7 4
4.1 4.2 4.3 4.4 4.5 4.6 4.7 4.8 4.9 4.10 4.11 4.12 4.13 5
5.1 5.2 5.3 5.4
Composite polymer electrolytes for electrochemical devices F. Alloin and C. Iojoiu, LEPMI CNRS, Grenoble Institute of Technology, France Introduction Composite electrolytes for lithium batteries: introduction Solid polymer electrolytes Composite polymer electrolytes based on poly(ethylene oxide) and clays Composite polymer electrolytes based on poly(ethylene oxide) and non-ionic fillers Gel polymer electrolytes Composite electrolytes for proton exchange membrane fuel cells Composite polymer electrolytes based on metal oxides Hygroscopic solid inorganic proton conductor composite polymer electrolytes Self-humidifying composite electrolytes Future trends Sources of further information and advice References Lithium-doped hybrid polymer electrolytes V. de Zea Bermudez, University of Trás-os-Montes e Alto Douro, Portugal, and M. M. Silva, University of Minho, Portugal Introduction Ionic conductivity Thermal properties Electrochemical stability
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95 97 103 110 112 123 124
129
129 130 131 132 137 148 151 156 160 164 166 167 168 176
176 178 182 184
Contents
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5.5 5.6 5.7 5.8
Spectroscopic studies Electrochromic displays Conclusion References
185 207 209 210
6
Hybrid inorganic–organic polymer electrolytes V. di Noto, E. Negro and S. Lavina, University of Padova, Italy, and M. Vittadello, City University of New York, USA Introduction Fundamentals of polymer electrolytes Overview of hybrid inorganic–organic polymer electrolytes Methods The real component of the conductivity spectra in the framework of the jump relaxation model and polymer segmental motion Conclusions Acknowledgements References
219
6.1 6.2 6.3 6.4 6.5
6.6 6.7 6.8 7
7.1 7.2 7.3 7.4 7.5 7.6 7.7 7.8 7.9 7.10 7.11 7.12 7.13 8
8.1
Using nuclear magnetic resonance spectroscopy in polymer electrolyte research S. Abbrent, University of South Bohemia, Czech Republic, and S. Greenbaum, Hunter College of the City University of New York, USA Introduction Nuclei possibility Liquid state nuclear magnetic resonance Solid state nuclear magnetic resonance Relaxation processes Diffusion measurements Magic angle spinning Double resonance experiments Two-dimensional methods Exchange nuclear magnetic resonance Electrophoretic nuclear magnetic resonance Conclusions References Molecular dynamics simulations of Li ion and H-conduction in polymer electrolytes D. Brandell, Uppsala University, Sweden Introduction
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219 220 221 254
266 272 273 273
278
278 279 281 282 291 295 299 300 304 305 307 308 309
314 314
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Contents
8.2 8.3 8.4
Computational chemistry The molecular dynamics methodology Li+-conducting poly(ethylene oxide)-based electrolytes for batteries Polymer electrolytes for fuel cells: perfluorosulphonic acid systems Conclusions and future trends References
8.5 8.6 8.7 9
9.1 9.2 9.3 9.4
9.5 9.6 9.7 9.8
Characterisation and modelling of multivalent polymer electrolytes M. J. C. Plancha, Laboratório Nacional de Energia e Geologia, Portugal Introduction Polymer–complex formation Ionic transport properties Morphological and crystallographic structures: characteristics and influence on ionic transport properties Ionic association: influence on ionic transport properties Phase diagrams: crystallinity and conductivity Conclusions References
Part II Applications 10
10.1 10.2 10.3 10.4 10.5 10.6 10.7 10.8 10.9
Polymer electrolytes for dye-sensitized solar cells J. N. de Freitas, J. E. Benedetti, F. S. Freitas, A. F. Nogueira and M. A. de Paoli, University of Campinas – UNICAMP, Brazil Introduction Polymer electrolytes Plasticized and gel polymer electrolytes Additives in the polymer electrolytes Stability of polymer electrolyte-based dye-sensitized solar cells Up-scaling: towards commercialization of polymer electrolyte-based dye-sensitized solar cells Conclusions and future trends Acknowledgements References © Woodhead Publishing Limited, 2010
315 316 325 329 336 337
340
340 342 344
358 365 367 373 373
379
381
381 384 390 402 417 421 423 424 424
Contents 11
11.1 11.2 11.3 11.4 11.5 11.6 11.7 11.8 11.9 11.10 11.11 11.12 12
12.1 12.2 12.3 12.4 12.5 12.6 12.7 12.8 13
13.1 13.2 13.3 13.4 13.5 14
14.1
ix
Solid polymer electrolytes for supercapacitors A. B. Samui and P. Sivaraman, Naval Materials Research Laboratory, India Introduction Solid electrolytes Conduction in solid electrolytes Solid electrolytes in supercapacitors Conducting polymer electrodes Activated carbon electrodes Cation exchange membrane-based supercapacitors Current research activities Applications Conclusions List of abbreviations References
431
Polymer electrolytes for electrochromic devices X. Fu, College of Chemistry and Chemical Engineering Southwest University, P.R. China Introduction Electrochromic effect and electrochromic devices Electrolytes for electrochromic devices Polymer matrix Classification of polymer electrolytes Proton-conducting polymer electrolytes and alkaline polymer electrolytes New type of polymer electrolyte References
471
Hyperbranched polymer electrolytes for high temperature fuel cells T. Itoh, Mie University, Japan Introduction Hyperbranched polymer electrolytes with a sulfonic acid group at the periphery Hyperbranched polymer electrolyte with a phosphonic acid group at the periphery Conclusions References Polymer electrolytes as solid solvents and their applications L. Ye and Z. Feng, Beijing Institute of Technology, China Introduction © Woodhead Publishing Limited, 2010
431 432 433 439 442 447 451 456 463 463 464 466
471 472 474 475 477 500 506 513
524 524 525 537 547 548
550 550
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Contents
14.2 14.3
Structure of lithium ion battery Advantages of polymer electrolytes in lithium ion batteries Main properties of polymer electrolytes Solid polymer electrolytes applied in lithium ion batteries Gel polymer electrolytes in lithium ion batteries Composite polymer electrolytes in lithium ion batteries Polymer electrolytes in other battery types Conclusions Acknowledgments References
14.4 14.5 14.6 14.7 14.8 14.9 14.10 14.11 15
15.1 15.2 15.3 15.4 15.5 15.6 15.7
Hybrid polymer electrolytes for electrochemical devices F. L. de Souza, Federal University of ABC, Brazil, and E. R. Leite, Federal University of São Carlos, Brazil Introduction Physicochemical properties of hybrid polyelectrolytes General discussion Applications Conclusions Acknowledgments References
Index
552 552 554 557 566 570 572 577 578 578
583
583 587 595 596 598 599 599
603
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Contributor contact details
(* = main contact)
Chapter 2
Editors and Chapter 1 Professor César A. C. Sequeira Department of Chemical and Biological Engineering Instituto Superior Técnico Technical University of Lisbon Av. Rovisco Pais 1049-001 Lisboa Portugal E-mail:
[email protected] Dr Diogo M. F. Santos Department of Chemical and Biological Engineering Instituto Superior Técnico Technical University of Lisbon Av. Rovisco Pais 1049-001 Lisboa Portugal E-mail:
[email protected]
J. S. Syzdek Warsaw University of Technology Faculty of Chemistry Inorganic Chemistry and Solid State Technology Division Polymer Ionics Research Group ul. Noakowskiego 3 00664 Warszawa Poland E-mail:
[email protected] and Université de Picardie Jules Verne Laboratoire de Réactivité et de Chimie des Solides 33 Rue St Leu 80039 Amiens Cedex France
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Contributor contact details
Chapter 3 A. Pawlicka* Instituto de Química de São Carlos Universidade de São Paulo PO Box 780 13560-970 São Carlos SP, Brazil E-mail:
[email protected] J. P. Donoso Instituto de Física de São Carlos Universidade de São Paulo PO Box 369 13560-970 São Carlos SP, Brazil E-mail:
[email protected]
M. M. Silva Department of Chemistry/Center of Chemistry University of Minho 4710-057 Braga Portugal E-mail:
[email protected]
Chapter 6 V. Di Noto,* E. Negro and S. Lavina Department of Chemical Sciences University of Padova Via Marzolo 1 I-35131 Padova (PD) Italy E-mail:
[email protected]
Chapter 4 F. Alloin* and C. Iojoiu Laboratoire d’Electrochimie et de Physico-chimie des Matériaux et des Interfaces (LEPMI) CNRS-Grenoble-INP-UJF Domaine Universitaire BP 75 F38402 St Martin d’Hères Cedex France E-mail:
[email protected];
[email protected]
Chapter 5
V. Di Noto CNR-ISTM Via Marzolo 1 I-35131 Padova (PD) Italy M. Vittadello PECS Department City University of New York/ Medgar Evers College 1650 Bedford Avenue Brooklyn, NY 11225 USA
V. de Zea Bermudez* Department of Chemistry and CQ-VR University of Trás-os-Montes e Alto Douro 5000-911 Vila Real Portugal E-mail:
[email protected]
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Contributor contact details
xiii
Chapter 7
Chapter 10
S. Abbrent* University of South Bohemia Branisˇovska 31 370 05 Cˇ eske Budeˇjovice Czech Republic E-mail:
[email protected]
J. N. de Freitas, J. E. Benedetti, F. S. Freitas, A. F. Nogueira* and M. A. De Paoli* Institute of Chemistry University of Campinas – UNICAMP PO Box 6154, 13083-970 Campinas SP Brazil E-mail:
[email protected];
[email protected]
S. Greenbaum Hunter College of the City University of New York 695 Park Avenue New York NY 10065 USA E-mail: steve.greenbaum@hunter. cuny.edu
Chapter 8 D. Brandell Department of Materials Chemistry Uppsala University PO Box 538 SE-751 21 Uppsala Sweden E-mail: daniel.brandell@mkem. uu.se
Chapter 9 M. J. C. Plancha Unidade de Pilhas de Combustível e de Hidrogénio Laboratório Nacional de Energia e Geologia, I.P. Estrada do Paço do Lumiar, 22 1649-038 Lisboa Portugal E-mail:
[email protected]
Chapter 11 A. B. Samui* and P. Sivaraman Polymer Division Naval Materials Research Laboratory Shil Badlapur Road, Ambernath (E) Thane, Maharashtra-421506 India E-mail:
[email protected]
Chapter 12 X. Fu College of Chemistry and Chemical Engineering Research Institute of Applied Chemistry Southwest University The Key Laboratory of Applied Chemistry of Chongqing Municipality Chongqing 400715 P.R. China E-mail:
[email protected]
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Contributor contact details
Chapter 13
Chapter 15
T. Itoh Division of Chemistry for Materials Graduate School of Engineering Mie University 1577 Kurima Machiya-cho Tsu City Mie 514-8507 Japan E-mail:
[email protected]
F. L. de Souza* Grupo de Materiais e Métodos Avançados – GMAv Centro de Ciências Naturais e Humanas (CCNH) Federal University of ABC Santa Adélia, 166, 09210-170 – Santo André São Paulo Brazil E-mail:
[email protected]
Chapter 14 L. Ye and Z. Feng* School of Materials Science and Engineering Beijing Institute of Technology No.5 Zhongguancun South Road, Haidian District Beijing 100081 China E-mail:
[email protected]
E. R. Leite Department of Chemistry Federal University of São Carlos, Rod. Washington Luiz, Km 235, 13565-905 São Carlos São Paulo Brazil
© Woodhead Publishing Limited, 2010
Preface
Polymers represent a class of materials currently in use in a broad variety of applications. The first meaningful application occurred in 1935 when Wallace E. Carothers invented nylon, which quickly found application in stockings, fabrics, ropes, combs, tyres and many household items. Along with metals, semiconductors, and ceramics, polymers serve as one of the cornerstones of materials science and engineering. A recent study by the National Research Council (UK) outlined some of the primary applications of polymers. Already, plastics such as styrofoam (polystyrene) are actively competing with paper in the packaging field. High strength plastics, such as polycarbonates and acrylonitrile–butadiene– styrene (ABS), serve as alternatives for house walls and roofing materials. Kitchen surfaces employ such plastics as polyethylene terephthalate (PET). Polymers are used in dentistry and in a variety of implantable drug-delivery mechanisms. In electronics, the high conductivity of certain classes of polymeric materials makes them attractive alternatives to conventional metallic wiring. They are used as photoresists in microlithography processes and are utilised as a packaging material for microelectronic devices. They have been used in optoelectronic devices as non-linear optical elements and as lightemitting diodes. They serve as sensors for chemical and biochemical agents. Polymers may act as hosts for ions in much the same way that liquid solvents do. Suppose that a salt is introduced into a polymer: when the solvation energy of the ions in the polymer host exceeds the ionisation energy of the salt, the individual ions separate and attach themselves to Lewis base sites on the polymer. Conductivity is possible above the glass transition temperature, where the polymer molecules are free to move. The free volume also provides room for the ions to move. The ionic conductivity possesses a diffuse liquid-like behaviour in addition to the hopping-type behaviour characteristic of an electron in a disordered solid. One way of thinking about how the conductivity is brought about relies on the dynamic-percolation model in which one envisages sites on a lattice interconnected by a set of randomly placed conductors, whose distribution can be changed continuously. xv © Woodhead Publishing Limited, 2010
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Preface
One of the goals of this book is to review the types, structure, geometry, topology, conductivity and other fundamental aspects that are unique to these ionic conductors, which were first introduced by P.V. Wright in 1975 and proposed as material of interest for the development of electrochemical devices by M. Armand in 1978. By using polymer electrolytes it is possible to design solid state batteries and related devices. The second goal of this book discusses a set of current applications of polymer electrolytes, namely high energy density batteries, fuel cells, supercapacitors, electrochromic devices, etc. An attempt is made to relate the applications to the underlying science described in the first part of the book. The work is the result of the joint efforts of many colleagues who contributed to it or helped with the editing process. First of all we would like to express our sincere thanks to all of the authors who supplied their high quality contributions within the short time frame available. This turned out to be a challenge for authors, editors and publishers. Furthermore, thanks are due to our families from whom we have stolen a number of weekends. César Sequeira Diogo Santos Lisbon
© Woodhead Publishing Limited, 2010
1 Introduction to polymer electrolyte materials C. A. C. S E Q U E I R A and D. M. F. S A N T O S, Technical University of Lisbon, Portugal
Abstract: Polymer electrolytes are electrolytic materials that offer many advantages in the area of large, high energy density batteries for electric propulsion and in fuel cells for electric vehicle or stationary applications. At the other end of the spectrum, polymer electrolytes are of interest for small, portable electronic devices where the battery represents a significant proportion of the device’s size and weight. This chapter focuses on relevant structural, physical and electrical properties of these materials and on the related sectors where there has been considerable industrial input. Key words: polymer electrolyte classes, polymer electrolyte properties, batteries, fuel cells and capacitors, photoelectro chemical devices, electrochromic applications.
1.1
Introduction
Polymers are at the forefront of the materials revolution; their uses are expected to grow at a faster rate (~15% per year) than that of any other structural compound. This trend may even be accentuated as polymers with new functionalities appear. We shall discuss here polymers having electrolyte properties for which important applications are foreseen. Ionic conductivity is usually restricted to salts in the molten state or their solutions in polar solvents. Despite their ubiquity, the understanding of such complex media is far from complete. On the other hand, a few remarkable solids with high ion mobility have been discovered, ranging from PbF2 by Faraday (1839) to recent newcomers including β-alumina and sulphide glasses. In such materials, one of the sublattices is in a quasi-molten state while still held by a rigid framework. Clays, zeolites, and natural and synthetic ion exchange resins may be considered as intermediate materials between solutions and the above-mentioned solid electrolytes. This is particularly true in the sense that a fixed network bears immobilised electrical charges while the discrete counter-ions are mobile, but only when the coulombic interactions are screened by a polar solvent. Water or alcohols are usually required as both anion (hydrogen bonding) and cation (electron pairs) solvation is needed (Gray, 1997). The motion of ions in polymeric 3 © Woodhead Publishing Limited, 2010
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Polymer electrolytes
matrices in the absence of a solvent is a relatively new phenomenon whose existence and importance have been recognised only in the past decades (Armand, 1986; Bruce, 1995a; Gray, 1997; Linford, 1987; Lipkowski and Ross, 1994; MacCallum and Vincent, 1987; Skotheim, 1988; Takahashi, 1989; Vincent, 1987a; Wright, 1975) and corresponds to a different concept; the macromolecule itself acts as a solvent for a salt which becomes partially dissociated in the matrix, leading to electrolyte behaviour. No small molecules are required for the conducting process, though the polymer–salt complex is routinely made from a solution of both constituents. An interesting parallel can be drawn between electronically and ionically conducting polymers: the latter (colourless) donate electron pairs from heteroatoms included in the macromolecular array from ‘s’ orbitals to Lewis acids (cations), while the former (highly absorbing) release single electrons from their conjugated ‘π’ system upon oxidative doping, with the formation of polarons (radical cations) possibly as dimers (bipolarons). Polymer electrolytes are to solutions what redox polymers (polypyrrole, polythiophene) are to metals. As a rule, electronic conductivity is best in spatially regular systems (crystals) while ionic conductivity prefers disordered matter. The ionic conductivity of polymer electrolytes is typically 100 to 1000 times less than exhibited by a liquid- or ceramic-based electrolyte. Although higher conductivities are preferable, and indeed a great deal of effort has gone into improving the bulk conductivity of polymer electrolytes over the years, 100-fold or 1000-fold increases are not essential, as a thin film electrochemical cell configuration can largely compensate for the lower values. Many polymer electrolyte materials will exhibit to a greater or lesser extent the following properties: • • • • •
adequate ionic conductivity for practical purposes; low electronic conductivity; good mechanical properties; chemical, electrochemical and photochemical stability; ease of processing.
These critical attributes are necessary if the materials are to be considered as practical replacements for their liquid counterparts. In addition, their properties, particularly conductivity and transport properties, should be sufficiently practical to stimulate their development when compared with other highly conducting solid electrolyte materials. Since 1978, when Michel Armand first introduced polyether–alkali–metal salt complexes to the solid state community as potential materials for electrochemical devices, there has been an enormous amount of research carried out on these (particularly high molecular weight poly(ethylene oxide)–lithium salt) systems, to obtain
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Introduction to polymer electrolyte materials
5
a better understanding of their fundamental properties and to use this information for the development of new generations of polymer electrolyte materials that are commercially more attractive (Armand, 1983; Plancha et al., 1997a).
1.2
Categories of polymer electrolytes
Rather confusingly, the term polymer electrolyte is widely used in fuel cell technology to refer to ion-selective polymers which, when hydrated, have two intermingled phases – the solid polymer region and a liquid-like region. An example is DuPont’s Nafion which is a fluorinated polymer with sulphonate end groups. The liquid phase is sometimes described as composed of micelles which are channel-like regions filled with liquid. Electrostatic regions in the ionic structure of the polymer backbone, with the mobile charges in water, help to form these regions which contain a liquid phase. There have been attempts to develop similar single ion conductive polymers for lithium batteries but, as yet, all materials reported have a resistivity which is too high for practical use. Finally, the term polymer electrolyte is also used for compositions which are essentially liquids adsorbed into a polymer. An example is found in the Bellcore system (Gozdz et al., 1994). US Patent 5, 296, 318 describes a copolymer of hexafluoride propylene and vinylidene 1,2-difluoride which is wetted with a liquid electrolyte. The term polymer electrolyte has, therefore, come to mean any polymerbased structure with significant ionic conductivity. These are found with a wide range of solid-like character. The solid character of polymers is, in general, related to the molecular weight of the polymer. Low molecular weight polymers are often liquids so the range can be from liquids to very hard and rigid materials. Some polymers can organise at the molecular level in such a fashion as to be crystalline. Since conductivity comes about through molecular motion in the structure, crystalline polymers have low conductivity and are not options for batteries. The ‘dry’ polymer electrolyte is a single phase, non-crystalline material containing dissolved salt, and where the ions of the salt are mobile. Plasticised polymers, a term from the polymer industry, are single phase and contain organic additives which have the effect of softening the polymer. These, in general, have higher conductivity than the dry polymers because of greater freedom for molecular motion. Solventdoped polymer electrolytes may be either single phase or two-phase depending on the degree of doping and the molecular make-up. The twophase materials are frequently described as gels. Both anions and cations are mobile at the molecular level. Finally, ion-selective, two-phase polymer electrolytes are in use in fuel cell technology but have not found application in lithium batteries (Gray, 1997).
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Polymer electrolytes
Among the different types of polymer electrolytes, the (liquid) solventfree system, where the ionically conducting phase is formed by dissolving salts in a high molecular weight polar polymer matrix, constitutes the ‘polymer electrolyte’ in its original sense. This category is the main subject of this chapter, due to its intensive development for lithium secondary batteries, electric propulsion applications, etc. Another category of primary importance is the proton-conducting polymer electrolyte used as the electrolyte in proton-exchange membrane (PEM) fuel cells, also known as solid polymer electrolyte fuel cells (SPEFC). These two categories will be further considered in the following sub-sections.
1.2.1 Lithium polymer electrolytes The history of lithium polymer electrolytes begins in 1973 with work by Fenton et al. (1973) at the University of Sussex (UK), and continues with several works by Armand (1986) at the University of Grenoble. Subsequently, two separate groups – a Harwell-led consortium of companies and a Hydro Quebec-Elf Aquitaine joint venture – set out to explore commercial possibilities of polymer electrolytes for lithium batteries (Araújo and Sequeira, 1994; North et al., 1982; Patrick et al., 1987; Sequeira and Hooper, 1983a). The first model to explain conductivity in polymer electrolytes emphasised the helical structure of poly(ethylene oxide) (PEO) and postulated the lithium cation either moved along the axis of the helix or along the outside of the strand. This model is based on ordered structures. It is not able to explain why the conductivity of PEO (the material can be described as solid if the molecular weight is >100 000) sharply changes at temperatures around 65 °C. By the mid-1980s it was recognised that most of the conduction occurred in the amorphous phase and the reason for low conductivity below about 65 °C was the polymer crystallinity. Subsequently, efforts were applied to reducing the crystallisation temperature of PEO (and other polymers). One method is to attach PEO strands to a polymer backbone. The relatively short strands are unable to self-align into a crystalline structure, thus extending the amorphous phase to much lower temperatures. Another option is to form copolymers which are not able to self-align as readily as the PEO strands. These types of material are what Megahed and Scrosati (1994) call the second generation of polymer electrolytes. Several solid polymer electrolytes (SPEs) based on copolymers have been prepared and complexed with lithium salts to prepare polymer electrolytes for lithium batteries. Incorporation of 10 to 20% poly(propylene oxide) (PPO) units in PEO depresses considerably the crystallinity of the materials. In addition to copolymer containing polyether grafted polyether, polysiloxane and polyphosphazene backbones have been used. Polymer
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Introduction to polymer electrolyte materials
7
electrolytes prepared with lithium salts and copolymers have shown enhanced transport properties at room temperature but conductivities and transport numbers are not sufficiently high to be attractive for practical lithium batteries (Fauteux et al., 1995). By the mid-1980s, measurements of transport numbers were showing that the anion was, at least, carrying a current equal to the cation and, in many measurements, a current greater than that carried by the cation. This observation made the original conductivity model inadequate. Molecular models have now been developed that describe motion of the ions based on oscillations of the polymer strands. And the free volume model for conductivity has been adopted from polymer chemistry. It is found that simple salts, such as LiCl, do not provide high conductivity. In general, the bulkier the anion, the higher is the conductivity. For example, Armand (1994) synthesised lithium bis(trifluoromethylsulphonyl)imide, LiN(CF3SO2)2, and observed high conductivity when dissolved in PEO-type polymer electrolytes. The high conductivity is the result of a more ‘plasticised’ polymer. Such an approach showed promising results with LiN(CF3SO2)2, through an increase in amorphous phase in the polymer electrolyte and an increase in conductivity. Therefore, new salts with perfluorosulphonate-polyether based single ion conductors have been synthesised. These salts have been incorporated into solvating polymers by co-crosslinking or copolymerisation. The resulting polymer electrolytes have been characterised by conductivity measurements. The room temperature conductivities of lithium polymer electrolytes are slightly better than the conventional polymer electrolytes but they are not high enough for practical applications in lithium batteries. Both the lithium cation and the anion of the dissolved lithium salt are mobile in the amorphous polymer. As a result, polarisation effects are encountered in battery operation. Efforts (e.g. Shriver at Northwestern University) to prepare single mobile ion electrolytes by building the anion into the polymer structure have not produced practical electrolytes. In all cases, the obtained conductivity has been at least an order of magnitude less than the comparable double ion structure (MacCallum and Vincent, 1987). The concentration of salt in the polymer matrix is higher than generally found in liquid solvents. It happens that the oxygen atom in the ethylene oxide backbone of PEO has just the right charge density to solvate the lithium cation but not to bind the lithium so tightly that it is immobile. On average, the lithium cation interacts with four ether oxygen atoms of the polymer (but not with two sequential ether oxygen atoms in the same strand). It has been observed that ion pairing sets in when the salt concentration exceeds a ratio of one lithium cation to eight ether oxygen atoms of the PEO strands and that ion aggregates appear at concentrations of one
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Polymer electrolytes
cation to four ether oxygen atoms. A maximum in conductivity occurs at concentrations around one lithium cation to eight ether oxygen atoms. The third generation of polymer electrolytes is variously described as containing a plasticiser, as dilute solvent or as gelled (Andrieu et al., 1995; Benrabah et al., 1995). Each of these terms may have had precise meaning at one time or another, but recently the terms seem to be used indiscriminately. Towards 1990, a group at the Engineering Research Laboratories in Denmark, with a group at the Mead Paper Company in the US (in what came to be known as the MHB venture) developed a solvent-diluted electrolyte which was polymerised in situ (Lee et al., 1989). This electrolyte is characterised by the incorporation of a liquid matter, such as propylene carbonate in a polymer matrix that is formed by crosslinking a monomer after mixing in the liquid phase. Following this, many methods and recipes for solvent-doped or ‘plasticised’ polymer electrolytes have been developed. (The term plasticised frequently means that the salt anion is an organic which, when dissolved, makes the polymer soft. But it is also used to refer to softening the polymer by addition of an organic liquid. In other ways, the hardness of the electrolyte depends on the molecular size and constitution of the polymer, so either dry or solvent-doped polymer electrolytes can range from essentially liquid-like to hard-brittle materials.) Surprisingly, up to about 80 wt% of some solvents can be incorporated into the matrix and yet it has the appearance of a dry solid. Other compositions have a texture more that of a gel. Even in the solid matrix, if the solvent has liquid-like character this can be observed by thermogravimetric analysis. The principal feature of the analysis is evaporation of the solvent at the temperature of the bulk liquid. Some compositions, however, do not undergo mass loss until temperatures are reached well above the boiling point of the bulk solvent. These compositions appear to be truly homogeneous ‘dry’ materials. Examples of plasticised polymer electrolyte systems with room temperature conductivities in the range of 2.6 × 10−3 S cm−1 are the polyacrylonitrile (PAN), polyvinylpyrrolidone (PVP), poly(methyl methacrylate) (PMMA) and poly(ethylene oxide) diacrylate (PEODA) plasticised with a low molecular weight solvent such as propylene carbonate, gamma butyrolactone or a high molecular weight solvent such as tetraethylene glycol dimethyl ether (tetraglyme) (Abraham and Alamgir, 1990; Cazzanelli et al., 1995; Huq et al., 1992). The main disadvantage with these polymer electrolytes is the reactivity of these solvents plasticised with lithium metal, accompanied by passivation and short cycle-life in lithium batteries. Also, for room temperature operation, molten salts containing 1-butylpyridinium chloride or bromide and aluminium chloride have been used for high conductivity polymer electrolytes. Poly(l-butyl-4-vinyl pyridinium) halides are known to be compati-
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ble with the molten salts to form polymer electrolytes. The addition of a small amount of the polypyridinium salts provides viscoelastic rubbery properties to the polymer electrolytes. The conductivity of the polymer electrolyte is influenced by the electrolyte composition and temperature (Watanabe et al., 1995). However, these electrolytes containing chloride ions, and particularly aluminium chloride are too corrosive and their processability into thin film form with mechanical stability is yet to be shown. They are not compatible with lithium anode and possibly not with many potential cathode materials to be useful for practical battery operation. Also in 1995, three superacid salts of lithium of the type CnF2n+1SO3Li where n = 4, 8 and 10 have been prepared and used in the preparation of SPEs with PEO (Nagasubramanian et al., 1995). The room temperature conductivity data show that these electrolytes have low conductivities. The conductivity of the polymer electrolytes PEO/C4F9SO3Li and PEO/C10F21SO3Li is of about 10−6 S cm−1, whereas the conductivity of PEO/C8F17SO3Li is < 10−6 S cm−1. These conductivity values are much lower than the conductivity of the polymer electrolyte with lithium bis(trifluoromethylsulphonyl)imide, PEO/(CF3SO2)2NLi, with a conductivity of 10−5 S cm−1. The increasing need for novel electrolytes that have high conductivity, high lithium ion transport number, a wide electrochemical stability window, and are suitable for lithium batteries with safe operation (Armand, 1994) is stimulating the third generation materials to be amorphous to temperatures as low as −40 °C and have conductivity at room temperature of the order of 10−3 S cm−1, approaching that of liquid electrolytes. Chiang et al. (1987) focused their attention on this need and produced interpenetrating polymer networks (IPNs) which maintained good electrical conductivity at room temperature. They synthesised IPNs from PEO– LiClO4 complexes and epoxy (bisphenol A diglycidyl ether) crosslinked with poly(propylene oxide) (PPO) triamine (Texaco Jeffamine T-403). The IPNs could be described as epoxy–PPO networks plasticised with PEG400. In this system, the epoxy and Jeffamine intercombined to form a solid matrix upon which PEO was held in an amorphous conformation at room temperature, as well as at higher temperatures. An analogy that could be drawn to this system is that of a wet sponge in which the water corresponded to the conducting liquid PEO–LiClO4 complex and the sponge corresponded to the solid support matrix of the epoxy and Jeffamine. Quantitatively, it was found that the introduction of the epoxy increased the mechanical properties of the polymer but decreased its conductive properties. When the epoxy content was greater than 20 wt% (w/o) (of the total mixture of lithium perchlorate (LiClO4) and PEO), free-standing films were formed and their mechanical properties improved as the percentage of epoxy increased. The conductivity, however, of the IPN remained nearly
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constant from pure PEO–LiClO4 to about 30 w/o epoxy. A further increase in the epoxy wt% led to an exponential decrease in the conductivity. Chiang et al. (1987) therefore concluded from their results that improved properties for a solid electrolyte were obtained with a composition of about 30 w/o epoxy in the IPN system. We have extended the study of Chiang et al. and had three main aims (Hudson and Sequeira, 1995): (i) to prepare ‘solid membrane films’ that could be properly tested – this involved finding suitable proportions of epoxy and Jeffamine which would form a suitable solid material; (ii) to investigate how the conductivity was affected by varying the amounts of LiClO4 present in the system; (iii) finally, to investigate the preparation and conductivity of an IPN in which the epoxy was crosslinked with tetraethylenepentamine rather than PPO triamine. This study showed that thin films (~0.015 cm) of IPNs could be prepared by a simple process, and have good mechanical and conductive properties, which are two of the essential requirements of a SPE. The ease of preparation and handling of these films would be vital if the IPNs fulfilled their commercial potential, if the large-scale mixing of different materials to form solutions of the IPNs would not cause the chemical engineer too many problems, and if the casting of the films would also be quite simple on a large scale. Moreover, the work showed that the conductivity of the IPN system was dependent on the concentration of the salt. The conductivity did not increase directly with increasing amounts of added salt because at very high salt concentrations neutral ion pairs were probably formed. Instead, there seemed to be an optimum value for the conductivity, which appeared at Table 1.1 The composition and room temperature conductivity of some solid polymer electrolytes (SPEs) and plasticised polymer films
Lithium salt PEO complex
Li : O ratio
Room temp. conductivity (S cm−1)
[CF3(CF2)3SO2]NLi : PEO [CF3(CF2)3SO2]NLi : PEO [CF3(CF2)3SO2]NLi : PEO [CF3(CF2)3SO2]NLi : PEO Lithium salt plasticised with PAN [CF3(CF2)3SO2NLi2] : PEO [CF3(CF2)3SO2NLi2] : PEO Dilithium salt plasticised with PAN C4F9SO2NLi2 plasticised with PAN C8F17SO2NLi2 : PEO C8F17SO2NLi2 : PEO LiClO4 : PEO in IPN LiClO4 : PEO in IPN
1:6 1:8 1 : 16 1 : 24 – 1:8 1 : 24 – – 1:8 1 : 16 1 : 88 1 : 176
5.31 3.24 3.0 4.77 1.02 1.79 3.03 8.06 1.1 4.24 4.37 6.10 1.80
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× × × × × × × × × × × × ×
10−6 10−6 10−6 10−6 10−4 10−6 10−6 10−5 10−5 10−7 10−7 10−5 10−6
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quite a low salt concentration where LiClO4 probably existed as mobile, free ions. The maximum and minimum conductivities at room temperature were obtained with (PEO)88:(LiClO4) and (PEO)176:(LiClO4), being 6.1 × 10−5 S cm−1 and 1.8 × 10−6 S cm−1, respectively. Compositions and conductivities of some SPE films and plasticised films that have been synthesised and characterised are listed in Table 1.1. Many more lithium polymer electrolytes have been developed for commercial applications (Araújo et al., 1996a; Cao et al., 2006; Dong and Wang, 2005; Hudson and Sequeira, 1993, 1994; Itoh et al., 2008; Kolonitsyn et al., 2005; Liao and Ye, 2004; Magistris et al., 2001; Nakagawa et al., 2003; Niedzicki et al., 2009; Niitani et al., 2005; Othman et al., 2007; Plancha et al., 1997b; Sequeira and Hooper, 1992; Shi et al., 2002; Tao et al., 2007; Wang et al., 2004). However, most of these lithium polymer electrolytes are still not good enough for lithium batteries and other electrochemical devices.
1.2.2 Proton polymer electrolytes The need to operate polymer electrolyte membrane (PEM) fuel cells at temperatures above 100 °C, where the amount of water in the membrane is restricted, has provided much of the motivation for understanding the mechanisms of proton conduction at low degrees of hydration. Although experiments have not provided any direct information, numerous theoretical investigations have begun to provide the basis for understanding the mechanisms of proton conduction in these nano-phase separated materials. Both the hydrated morphology and the nature of the confined water in the hydrophilic domains influence proton dissociation from the acidic sites (i.e. SO3H), transfer to the water environment, and transport through the membrane (Paddison, 2003). Hereinafter we discuss progress on proton membranes in SPEFCs. The Nafion membrane is an example of this type of membrane. Classically, the Nafion membranes are chemically synthesised in four steps according to the DuPont de Nemours process (Grott, 1978): (a) the reaction of tetrafluoroethylene (TFE) with SO3 to form the sulphone cycle; (b) the condensation of these products with sodium carbonate followed by copolymerisation with the TFE to form an insoluble resin; (c) the hydrolysis of this resin to form a persulphonic polymer; and (d) the chemical exchange of the counter ion Na+ with the proton, H+, in an appropriate electrolyte. This membrane was introduced by DuPont in 1966. The development of perfluorinated membranes by DuPont during the 1960s has played a vital role in electrochemical system applications. The Nafion family or perfluorinated ionomer membranes meet the requirements for several electrochemical systems (chloroalkaline, fuel cells and some other non-fuel cell applications). A lifetime of over 60 000 hours has
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Polymer electrolytes ( CF2 – CF2 )x ( CF – CF2 )y [OCF2CF]z – O(CF2)2 SO3H CF3
1.1 The general structure of Nafion membrane: x = 6 − 10 ; y = z = 1.
been achieved at 80 °C. Their general formula is given in Fig. 1.1 (Savadogo et al., 1995). The values of x, y and z in Fig. 1.1 are x = 6−10, y = z = 1. These values can be varied to produce materials with different equivalent weights (EWs) and pendant chain lengths, the equivalent weight (EW) being the number of grams of polymer per mole of fixed SO3 sites. The high EWs of the first Nafion membranes’ family limited their use in fuel cells and prompted the development of the Dow membrane. This membrane is structurally and morphologically similar to the Nafion membrane, but differs with respect to its EWs, which are typically in the 800 to 850 range, and have shorter size chains (z = 0 for Dow and z = 1 for Nafion) (Utracki and Weiss, 1989). The Dow membranes are a short side chain perfluorinated ionomer whereas the Nafion membranes are the long side chain perfluorinated ionomer. The specific conductance of 800 and 850 EW experimental membranes has been reported as 0.20 and 0.12 Ω−1 cm−1, respectively (Savadogo et al., 1995). It must be pointed out that the Dow monomer is more complicated to elaborate than the DuPont monomer. Therefore, the synthesis of the Dow epoxy is more complicated than that of the Nafion, which is a commercially available material (Savadogo et al., 1995). Testing the Dow membrane, during the period 1987–88, in Ballard fuel cells resulted in a dramatic increase in the SPEFC performance levels (Watkins, 1988). Since this success by Dow, DuPont has been actively developing their membranes with respect to durability and continuous improvement. They have increased power densities by further decreasing the equivalent weight, from 1100 EW (Nafion 117) to 1000 EW (Nafion 105), and the membrane thickness from 178 to 127 µm. Nafion 115 (127 µm) and Nafion 105 (127 µm) are some of the latest materials developed specially for fuel cell applications. Perfluorinated ionomer membranes have also been developed by Asahi Chemical and the Asahi Glass Company and commercialised as Aciplex S and Flemion, respectively. The general properties of the long side chain perfluorinated ionomer membranes (e.g. Nafion, Flemion, Aciplex) and the short side chain perfluorinated ionomer membranes (e.g. Dow) are shown in Table 1.2.
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Table 1.2 Properties of perfluorinated monomer membranes (e.g. Nafion, Flemion, Aciplex, Dow) EW range
800–1500 −1
Conductivity range (S cm ) Conductance range (S cm−2) Dimensional stability
Proven lifetime Thickness (µm)
0.20–0.05; for example: 1100 EW = 0.10 and 850 EW = 0.15 2–20; for example: 1100 EW (178 and 51 µm membrane) = 5 and 17, respectively 10–30% expansion X and Y direction from 50% RH liquid water (25 °C); for example: 1000 EW = 16% expansion >50 000 hours for Nafion, Flemion and Aciplex; >10 000 hours for Dow 250–50
EW stands for equivalent weight and RH stands for the relative humidity of the membranes.
Nowadays, the long side chain perfluorinated polymer electrolytes have proven to have a prolonged service life under electrolysis and electrosynthesis conditions and in fuel cell systems (Appleby and Yeager, 1986; Blomen and Mugerwa, 1983; Chapiro, 1962; Cohen et al., 1949; GuzmanGarcia et al., 1992; Hodgdon, 1968; Küver et al., 1994; La Conti, 1988; Prater, 1990; Riedinger and Faul, 1988; Samms et al., 1996; Savadogo and Roberge, 1997; Wang et al., 1996a). They are not efficient for high temperature (>150 °C) fuel cell systems because their conductivity decreases significantly due to dehydration. Although some efficiency has been gained by lowering the membrane EWs, the main improvements in fuel cell performance were achieved by simply thinning the membranes (Prater, 1990). The advantages gained with this simple strategy include lower membrane resistance, lower material utilisation (which obviously allows cost savings), and improved hydration of the membrane. However, there is a limit to the extent to which such membranes can be thinned because of difficulties with durability and reactant crossover. W.L. Gore and Associates, Inc. proposed to solve some of these difficulties using perfluorinated monomer membranes reinforced with woven polytetrafluoroethylene (PTFE). Perfluorinated membranes based on weak acid functions have been developed by Asahi Chemicals (Japan). Their EW is similar to that of the above membranes. The exact synthesis process of the membrane is not indicated. They have been developed for the chlor-alkali industry where they are used in bilayer membranes. In these membranes a strongly acidic membrane (–SO3H) (pKa < 1) is coated on one side, and the second side is coated with a thin layer of a weak-acid membrane (–COOH) (pKa = 3).
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Polymer electrolytes
Only the strongly acidic membrane is exposed to the anolyte so that acid may be added while the weak-acid layer is exposed to the catholyte to prevent sodium hydroxide transport from catholyte to anolyte. A comparison between the membranes containing –SO3H and –COOH ion exchange groups has shown that the water content and the ionic conductivity of the –SO3H-based membranes are higher than those of the –COOH-based membranes. Therefore, the latter are less suitable for fuel cell applications. Perfluorinated membranes reinforced with woven PTFE, such as Nafion 324 or Nafion 417, are used in many industrial electrochemical processes. Unfortunately, the relatively coarse weave of the woven PTFE reinforcement results in membranes that are much too thick for high electrochemical performance. Some non-woven PTFE/perfluorinated ionomer composite membranes have been formulated for ion transport studies and for investigation of other properties relevant to chlor-alkali and fuel cell applications. Advancements in materials and processing technology have resulted in the introduction of PTFE/perfluorinated ionomer composite membranes claimed by Gore and Associates, Inc. under the Gore Select trademark. Most types of ion exchange membranes contain some type of reinforcement because they are weak and tend to swell substantially as they incorporate a solvent onto their structure. This reinforcement can be a ‘macro reinforcement’ or a ‘micro reinforcement’. The reinforced membranes are, of course, more hydrophobic than the conventional Nafion membranes. It may be interesting to investigate if the low resistance is due to the high reinforcement of the membrane (which is very thin) or to its water transport ability. It is well established that, for all types of membranes, the low EW enhances the water transport ability. But it does not seem that the reinforcement has a beneficial effect on water transport. Up to now, the Nafion membrane and the Dow membrane have been most effective advanced membranes suitable for use in practical systems. The Dow membrane exhibits superior performance to that of the Nafion 117 membrane. The major disadvantage of these membranes is their high cost. This is due to the long preparation process and to the thickness of these membranes. To achieve direct chemical synthesis of Nafion membranes without these problems, novel methods of preparation were developed. It was found that the ionic conductivities of the membranes elaborated with Nafion, heteropolyacids and/or thiophene are higher than those with just Nafion. The ionic conductivities of these membranes fabricated with Nafion electrolyte and silicotungstic acid (STA), with and without thiophene, are 2.3 times higher than that of the Nafion 117 membrane. Their conductances are also higher than that of the Nafion 117 membrane. The water uptake of the Nafion, Dow, Gore Select and the membranes fabricated with Nafion electrolyte and STA are shown in Table 1.3. As it
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Table 1.3 Polymer electrolyte membranes water uptake Membrane
EW
Thickness (µm)
Water uptake (%)
Nafion 117 Nafion 112 Dev Dow Gore-Select Gore-Select NASTA NASTHI NASTATHI
1100 1100 800 1100 900 1100 1100 1100
200 60 100 20 12 410 120 150
34 34 56 32 43 48 32 30
may be seen, the water uptake of the Dow membrane is higher than that of the Nafion 117. This is related to the chemical composition of the membrane and not to its thickness. Effectively, Nafion 117 (200 µm) and Nafion 112 (60 µm), which have the same chemical composition but different thicknesses, exhibit the same water uptake. The water uptake of the Nafion solution and STA is 40% higher than that of Nafion. This may indicate that the STA may modify the membrane chemical properties. In the same way, membranes fabricated with Nafion 117 and STA (NASTA) enhance water transport. Similar results were obtained on other heteropolyacids. The improvements on the membranes’ conductivity, conductance and water transport were attributed only to the presence of STA. The preparation of the monomer α,β,β-trifluorostyrene (TFS) was first carried out by Cohen et al. (1949) who reported its total synthesis. Hodgdon (1968) firstly investigated the sulphonation of TFS with applications to structures and cells. Optimising the reaction conditions, Hodgdon was able to prepare a multiplicity of EW or different ion-exchange capacity (IEC) of both linear and crosslinked poly-TFS sulphonic acids. He showed that the extreme difficulty in sulphonating poly-TFS was caused by the betadirecting influence of the perfluorinated polyalkyl group attached to the aromatic ring (Hodgdon, 1968). It was claimed that the high stability of the fluorine atoms attached to the alkali carbon atoms impacts oxidative and thermal stability superior to that exhibited by conventional ion exchange polymers such as polystyrene sulphonic acid (La Conti, 1988). The limited use of these membranes in fuel cells may be related to their limited mechanical properties and/or chemical resistance in real fuel cell conditions. Based on the above work on sulphonated TFS membranes, a novel family of sulphonated copolymers incorporating TFS and a series of substituted TFS monomers provided the group of materials referred to as BAM3G (Ballard Advanced Materials 3rd Generation of membranes) (Savadogo et al., 1995; Savadogo and Roberge, 1997).
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The most significant earliest work carried out on the radiation grafting of polymers was achieved and reviewed by Chapiro (1962). The cation exchange membrane (CEM) was prepared by graft polymerising TFS onto the inactive fluorine-containing polymer film, and then sulphonating the grafted film in different patterns. The grafted polymer membranes were obtained at a graft rate of 10 to 50%. The membrane containing grafted polystyrene sulphonic acid and two low density polystyrene membranes (PTFE and a copolymer of TFE and hexafluoropropylene) have shown similar conductivities to those of the Nafion and Dow membranes (Savadogo and Roberge, 1997). Many reports (Savadogo et al., 1995) indicate that the oxidation stability of these grafted foils is very limited. This makes them appropriate for fuel cell applications; but only the membranes based on the PTFE backbone showed some significance as candidate materials. It is well established that the conductivity of the perfluorinated ionomer membranes currently used in PEM fuel cells depends on the presence of water to solvate the protons generated by the ionisation of the sulphonate acid groups. When such polymers are subjected to temperatures above 100 °C at atmospheric pressure, dehydration occurs, and conductivity decreases significantly (Samms et al., 1996). It has been shown that for a prototype direct methanol fuel cell (DMFC), the methanol crossing over the membrane was responsible for the decrease in cathode potential by at least 100 mV (Küver et al., 1994). From the investigation of the methanol diffusion in Nafion (Küver et al., 1994) it was concluded that methanol readily transports across perfluorosulphonic acid (PFSA) membranes, and the investigation of a few membrane systems for DMFCs was suggested. The low cost and excellent oxidation and thermal stability of phosphoric acid doped polybenzimidazole (PBI) prompted researchers at Case Western University (Samms et al., 1996; Wang et al., 1996a,b,c) to develop this membrane as a polymer electrolyte for DMFCs. After investigation of the thermal stability of PBI doped with phosphoric acid up to 600 °C, it was concluded that this membrane is adequate for use as PEM in a high temperature fuel cell. These studies may also be extended to other polymers like the polybenzimidazobenzophenanthrolines (PBIPAs) which exhibit excellent thermal and mechanical properties (Zhou and Lu, 1994). The sulphonated polyimide membranes are based on 4,4′-diaminobiphenyl2,2′-disulphonic acid (BDSA), 4,4′-oxydianiline (ODA), 4,4′-oxydiphthalic anhydride (ODPA) and 1,4,5,8-naphthalenetetracarboxylic dianhydride (NTCDA). The study of the surface properties of fluorinated polyimides exposed to vacuum ultraviolet (VUV) radiation and atomic oxygen may help us understand the stability of these polymers. It will also be interesting to apply this sulphonation method to some novel high molar mass polyetherimide, as reported by Hsio and Yang (1997).
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The poly(benzylsulphonic acid) siloxane (PBSS)-based membrane is based on organic–inorganic protonic polymer electrolyte (ormolyte–organically modified silane electrolyte) (Chambouillet et al., 1988; GauthierLuneau et al., 1992). The synthesis was carried out by the sol–gel process during different steps. It was based on the hydrolysis-condensation of the benzyltriethoxysilane precursor in a methanol solution containing water and trifluoromethanesulphonic acid (CF3SO3H) (Gauthier-Luneau et al., 1992). The final inorganic–organic ion-exchange ormolyte had good thermal stability up to 250 °C, and showed a conductivity of 1.6 × 10−2 Ω−1 cm−1 at room temperature (Gauthier-Luneau et al., 1992). It was claimed that these properties made them useful for DMFCs but, up to now, no fuel cell performance data were given. It would be interesting to study the methanol crossover in this membrane; its chemical stability in practical fuel cell systems should also be addressed. Besides the studies on anhydrous proton polymer electrolytes, various research groups have studied proton conductivity in systems in which aqueous solutions of H3PO4 or H2SO4 were used as additives to the polymeric matrix (Lassagues et al., 1992; Raducha et al., 1996; Wieczorek et al., 1994). It has been demonstrated that thin film proton conducting polymer gels based on polyacrylamide (PAM) matrix and agar with a H3PO4 additive exhibited reasonable mechanical strength and room temperature conductivities in the range of 10−3 to 10−2 S cm−1 (Wieczorek et al., 1995). The conductivities are higher than those obtained from many of the above non-fluorinated membranes. Differential scanning calorimetry (DSC) measurements indicate that these protonic gels are stable up to 120 °C. Their mechanical stability and conductivities are very dependent on the water content. Also, their water uptake is not well determined in various operation conditions. According to the available literature data, the conductivities of these systems are less than those of perfluorinated membranes at room temperature. It would be of interest to evaluate their performance in practical fuel cell systems. In particular, the stability of these proton conductors in actual practical systems must be addressed in detail. In summary, the modifications made to the Nafion membranes, the conceptual design of substitutes for PFSA materials and the modifications made to aromatic membranes to render them suitable for fuel cells and other electrochemical device applications constitute promising avenues for present and future research.
1.3
Structure and its implications
It is of maximum importance to understand the ion transport in polymer electrolytes. It is the basic principle governing the interactions between ions
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and their polymeric host (and the ions themselves). Of course, the short range dynamics of the polymer host almost invariably control the ratedetermining step for ionic motion in simple polymer electrolytes. Many efforts have been made over the years to address these issues, in particular through Raman, infrared, nuclear magnetic resonance (NMR) and extended X-ray absorption fine structure (EXAFS) studies of amorphous systems (Aihara et al., 2000; Bronstein et al., 2004; Forsyth et al., 2002, Frech and Manning, 1992; Glasse et al., 1992; Jin et al., 2000; Ouyang et al., 2007; Wintersgill et al., 1988). For a salt MX dissolved in a polymer host solvent, the formation of ion pairs, [MX]•, produces neutral species and the concentration of charge carriers drops. Larger aggregates may also exist (e.g. [M2X]+, [MX2]−), whose mobilities will be impaired by their size in comparison to free ions, and conductivity will again be adversely affected. The importance of ion association in a salt solution is largely dependent on the dielectric constant, ε, of the solvent. As the salt concentration is increased, the inter-ionic distance decreases and ion–ion interactions become progressively more significant. The onset of ion-pair formation occurs at lower ion concentrations in solvents of low ε. As ε ~ 5–10 for polyethers (cf. 78.5 for water), extensive ion–ion interactions are expected to be favourable. When a salt is dissolved in a polymer matrix, the conductivity of that polymer increases due to the concentration of charge carriers. Moreover, as the salt concentration, for example for a polyether, is increased above ~0.1 mol dm−3 (M:O ratio of ~ 1:100 to 1:50), the conductivity is found to reach a maximum and then falls. Ion transport is closely coupled to the segmental motion of the polymer. Cations moving between coordinating sites must adopt a transient state, where they are coordinated by both sites, before the motion of the polymer breaks one coordinating bond. The fall in conductivity as the salt concentration increases is thought to be due to the introduction of an ever-increasing number of the transient crosslinks in the system, which causes a reduction in chain mobility. Ion aggregation can still contribute to the observed decrease in conductivity as, if aggregates do exist, they are likely to have retarded diffusion rates, making an additional contribution to the fall in conductivity. Infrared and Raman spectroscopy have been used to study both interactions between the ions and host polymer and those between cations and anions (MacCallum and Vincent, 1989). This relies on monitoring charges in the vibrational modes of anions, namely the trifluoromethanesulphonate (CF3SO3−) and perchlorate (ClO4−) anions (MacCallum and Vincent, 1989). Variation in vibrational frequency and the Raman spectral line width can provide information about the influence of the environment on the molecular ions. Changes in dipole moment arising from an ion-pair vibration are Raman inactive but strong infrared absorptions, usually in the far infrared,
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are expected. In addition, Raman scattering studies have been used to quantify the degree of ion association and prove its variations with temperature and salt concentration (Ferry et al., 1995; Gray, 1991; Schantz and Torell, 1993). Raman and Fourier transform infrared (FTIR) spectroscopy studies have been carried out in a wide range of metal salt complexes. Many trends, such as increased pair/cluster formation with increased temperature or salt concentration, have been noted (Bernson and Lindgren, 1993; Wendsjö et al., 1992). As a result of the many studies carried out, a wider variety of configurations is possible in the complex systems, and complicated temperature dependence for solvation is expected (Andreev and Bruce, 2000; Papke et al., 1981a; Staunton et al., 2005; Thomson et al., 1996). Ion pairs or aggregates may be defined as contact or solvent separated species. The former is produced when a molecule of the first-neighbour solvent shell around an ion is replaced by an oppositely charged ion. The latter results when there is an overlap of solvent cospheres, linking a pair of ions to form a mechanical unit. The ion centre-to-centre distance is therefore greater in this case, not needing to be limited to one interlinking solvent molecule. In general, it is very easy for contact ion pairs to form in polymer electrolytes, even at low salt concentrations (Bruce and Gray, 1995). While the stoichiometric crystalline complexes do not contribute significantly to the conductivity of the polymer electrolyte, their structure provides valuable information on ion–polymer and ion–ion interactions. There is new evidence from infrared and Raman spectroscopy, EXAFS, thermodynamic and other studies that the local environment in the melt is likely to be very similar to that in the crystal from which it was obtained. Therefore, the crystal structure could provide a useful insight into the structural make-up of the conducting amorphous phases. X-ray powder diffraction has been used extensively to establish the unit cell dimensions of polymer electrolyte crystalline complexes. Determination of crystal structures of polymer electrolytes is not so straightforward since single crystal X-ray diffraction is not suitable. In addition, the large unit cells and low symmetry of polymer electrolyte crystalline phases cause conditioning problems. Success in polymer electrolyte crystal structure determination has come through high quality powder X-ray data of polycrystalline powders. These data can be analysed using Rietveld profile refinement in which an unrefined starting model for the structure is fitted to the observed data by a non-linear least squares technique. Alternatively they can be analysed by using powerful ab initio methods, where no starting model is required. To date, many structures have been elucidated in this way (Andreev et al., 1996; Baskaran et al., 2007; Bruce, 1995b; Chatani and Okamura, 1987; Frech et al., 1997; Fullerton-Shirey and Marenas, 2009; Henderson et al., 2003; Lightfoot et al., 1993; Martin-Litas et al., 2002; Murakami et al., 2002).
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There are many aspects of the structural features of polymer electrolytes that cannot yet be fully resolved experimentally and significant contributions could be made by undertaking computer simulations. Areas where simulation studies could be beneficial include the following: •
•
•
• •
Determination of the form that polymeric helices take-up above the melting point. Then, the relative importance of inter- and intra-chain crosslinks over a concentration range could be correlated with the glass temperature (Tg) and conductivity measurements. Crystal structure and vibrational spectroscopic data have only provided information on a single stoichiometric phase, that is, on short-range environments, so it would be useful to address this issue through computational methods. Spectroscopic data suggest that there are at least three environments for anions, i.e. different ion associated species, which vary in concentration at moderately high salt concentrations, and computer simulations could assign structures to these ionic clusters. Elucidating the role of various plasticisers in gel-type electrolytes. More structural information from simulation studies could play a role in designing gel electrolytes for optimum transport and dynamic properties.
In addition to building our structural knowledge, computer simulation studies can also play a key role in resolving ion transport processes. Several groups have carried out molecular dynamic simulations of polymer electrolytes. These simulations have been implemented in investigations of the polymer electrolytes’ structural properties and are revealing accurate descriptions of these materials and their potential applications (Capiglia et al., 2000; Catlow and Mills, 1995; Ennari et al., 2000; Gadjourova et al., 2001; Joo et al., 2002; Karlinsey et al., 2004; Neyertz et al., 1995; Payne et al., 1995; Ward et al., 1995; Zheng et al., 2000).
1.4
Conductivity measurements
The conductivity of a polymer electrolyte is an important parameter. This section addresses electrochemical conductivity techniques for the study of SPEs. The different types of conductivity are discussed, followed by an outline of the features, applicability and validity of direct current (DC) and alternating current (AC) conductivity measurements. Techniques for the identification of the individual species responsible for conduction are then briefly reviewed.
1.4.1 Contributions to conductivity The total conductivity, σ, is the sum of the electronic, σel, and ionic, σi, contributions © Woodhead Publishing Limited, 2010
Introduction to polymer electrolyte materials σ = σel + σi
21 [1.1]
Consider a polymer electrolyte of the type Pn:MX (where P is the structural repeat unit of the polymer chain and n is the stoichiometric ratio of structural repeat units to formula units of salt MX). Furthermore, assume that the valences of M and of X are plus one and minus one, respectively. An average value of the electrical conductivity can be obtained by measuring the AC conductivity of Pn:MX between two inert electrodes and with an inert gas passing over the sample. Alternatively, four probes can be used and the IR drop is measured between the two inner probes (see Fig. 1.2). However, both of these configurations leave the chemical potentials of M and X ill defined. It is known that the partial conductivities may be decisively affected by the M/X ratio, which will be established by fixing the chemical potentials. Because of this, it is preferable to fix the M/X ratio by fixing the chemical potential of one of the two constituents, as shown schematically in Fig. 1.3. Figure 1.3(a) involves symmetric electrodes of the parent metal M. These can fix the chemical potential of M. Figure 1.3(b) shows a setup using an alloy, MN, to fix the chemical potential of M at a value lower than that
Argon gas Pn : MX
Pt
Pt
IR drop
1.2 Schematic diagram of electrodes used to measure the AC conductivity of Pn:MX either by the two-probe or four-probe technique. The chemical potentials in the polymer electrolyte are not defined.
Argon gas (a)
Pn : MX
M
M
Argon gas (b)
M–N alloy
Pn : MX
Porous Pt (c) pX2
M–N alloy Porous Pt
Pn : MX
pX2
1.3 Schematic diagram of electrodes used to fix the M/X ratio in a solid polymer electrolyte, Pn:MX: (a) electrodes of the parent metal, M; (b) electrodes of an alloy, MN, in which N is inert; (c) electrodes of porous platinum with a coexisting partial pressure of X2 gas.
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Polymer electrolytes
depicted in Fig. 1.3(a). Here, N is an inert element. Figure 1.3(c) shows a diagram of electrodes to fix the chemical potential of M at its lowest value. In each of these cases, the electrodes are assumed to be reversible. Using such setups, one can obtain a well-defined value of the total conductivity σ as defined above. A useful electrolyte must behave not only as an ionic conductor but also as an electronic insulator. In practice, this means that the electronic conductivity must be so low (~10−9 S cm−1) that the cell would take a very long time, say four years, to self-discharge on standing. In general, the electronic conductivity of an electrolyte should be at least five orders of magnitude lower than the ionic conductivity, which means that the self-discharge current is below 10−5 of the normal current under load. Polymer electrolytes, unlike many other forms of solid electrolytes, have no difficulty in fulfilling this requirement. The electronic conductivity arises from current carried by electrons, σe, and/or electron holes, σh: σel = σe + σh
[1.2]
It is possible to discriminate between these two contributions by using techniques like the Wagner polarisation cell, which is discussed in Section 1.4.4. The disadvantage of such methods is that the test electrodes replace those that would be used in a practical device application and so the behaviour of the electrode–electrolyte interfaces is significantly different. In consequence, the predicted electronic conductivity may be some orders of magnitude higher or lower than that found when the electrolyte is incorporated in a working device. The ionic conductivity of a polymer electrolyte of the type Pn:MX would be expected to arise solely from cationic contributions from M+ and anionic contributions from X−. For polymer electrolytes where P is PEO, for which the relative molar mass of the repeat unit is 44, and n is typically about 4 to 10, the concentration of the ‘solution’ of salt in the polymer is about 2–5 M. This approaches the concentration found in molten salts. A feature of the behaviour of such concentrated systems is the preponderance of ion pairs and triple ions. The ionic transport then includes contributions from cations such as M2X+ and anions such as MX2−. Most techniques for determining the individual conductivities of different species in fact only discriminate between two classes: those that are reversible at a given non-blocking electrode and those that are not. If the test cell is configured so that the contribution of the cations to the total ionic conductivity can be determined, then the cationic transference number, τ+, can be obtained. This may be defined as: τ+ =
σ+ = σi
∑ (σ σi
+
)j
= ∑ ( t+ ) j
© Woodhead Publishing Limited, 2010
[1.3]
Introduction to polymer electrolyte materials
23
where (σ+)1, (σ+)2, etc., are the conductivities of the first, second, etc., cationically conducting species. The anionic transference number, τ−, is similarly defined, and is related to τ+ by τ− = 1 − τ+
[1.4]
The term transport number, t, is reserved for the contribution of an individual species. For the polymer electrolyte Pn:MX, there are separate transport numbers t (M+), t (M2X+) (and perhaps, if water is also present within the electrolyte, t (H+) for each of the individual mobile cation species, but it is easier to determine the sum of these t+, than to separate out the speciated values.
1.4.2 Direct current conductivity DC conditions using either blocking or non-blocking electrodes give rise to polarisation problems in conductivity measurements. In spite of these pitfalls, the total DC conductivity has routinely been used to characterise polymer electrolytes and a threshold value of approximately 10−5 S cm−1 has been used as the criterion for possible application purposes. A large number of materials reach such a conductivity value between room temperature and 100 °C (Aoki et al., 2006; Blonsky et al., 1986; Chiang et al., 1985; Dupon et al., 1984; Egashira et al., 2008; Fang et al., 1988; Geiculescu et al., 2004; Gray, 1990; Henderson, 2007; Marzantowicz et al., 2007; Xi et al., 2004; Y. Zhang et al., 2006). Several aspects govern the magnitude of the conductivity, namely the degree of crystallinity, the salt concentration, and the nature of the salt and polymer (Martins and Sequeira, 1990). Temperature and pressure have been used in a great number of studies as variables in investigations of these conductivity aspects. With blocking electrodes, there is neither a source nor a sink for mobile ions. The migration of the ions under the influence of the electric field leads to an enrichment of the mobile species in the electrolyte region adjacent to one electrode and depletion near the other electrode. The ionic motion is then opposed by a chemical potential gradient and when, after a short time, this has increased sufficiently to counterbalance the electric field, the migration stops. The cell is then said to be concentration polarised. Since the calculation of the cell conductance depends on the assumption of ohmic behaviour, which pertains only at the instant of initial application of the applied voltage, it is difficult to determine conductance accurately from DC measurements on test cells with blocking electrodes. The mechanism by which current initially flows across a test cell with blocking electrodes under the influence of an applied voltage can be visualised as follows. Neither the electronic charge carriers within the electrode nor the ionic charge carriers within the electrolyte can cross the electrode–
© Woodhead Publishing Limited, 2010
24
Polymer electrolytes Cg
Cdl (a)
Rb
Cdl
Cg
Rct
Rb (b)
1.4 (a) Equivalent circuit for a test cell with blocking electrodes. (b) Equivalent circuit for a test cell with non-blocking electrodes.
electrolyte interface. The applied voltage causes a build-up of negative charge on the electrode side of one of the interfaces. This induces a corresponding build-up of positively charged ions on the electrolyte side of the same interface. A similar ‘double layer’ of separated charges, but with the signs reversed, is set up at the other interface. The electric field that causes ionic motion within the electrolyte is that caused by the charge concentration gradient from the electrolyte side of the interface. The double layers each have infinite resistance to the direct passage of charge, but have a capacitance term that results from the layers of separated charge. The double layer can be visualised as a parallel-plate capacitor, and since ε = ε0εrA/d
[1.5]
(where ε0 is the permittivity of free space, εr is the relative permittivity, A is the area of the plates and d is the distance between them), the capacitance term, C, is relatively large (of the order of µF) since the separation, d, is very small. The equivalent circuit to the cell is displayed in Fig. 1.4(a). The so-called geometric capacitance, Cg, that appears in parallel is a consequence of the charges on the electrodes themselves. This term would be present even if the electrolyte were to be removed. Its magnitude is small compared with the double-layer capacitance, Cdl, because the intercharge separation is much greater. The electrolyte conductance, G, is the reciprocal of the electrolyte bulk resistance, Rb.
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Introduction to polymer electrolyte materials
25
In the case of measurements with non-blocking electrodes, there is both a source and a sink for the mobile ionic species. As in the case of blocking electrodes, however, DC measurements do not easily provide accurate values for Rb. This is because, although ions can cross the boundary, electrons and/or holes cannot. Since these are the conducting species within the electrodes, the charge-transporting process has to be transferred to and from electrons and ions at the two interfaces. The kinetics of this charge-transfer process is not infinitely facile and, consequently, there is a resistance, Rct, associated with the charge-transfer process. This results in a build-up of charge on either side of the interface, which produces an associated, and quite large, charge-transfer capacitance, in parallel with the resistance. The equivalent circuit for this situation is similar to that for blocking electrodes, but the double-layer capacitance is shunted by the charge-transfer resistance, as shown in Fig. 1.4(b). A DC measurement, therefore, does not give the value of Rb directly. It has been observed for both hydrated and dehydrated salts that isotherms of log (σ/S cm−1) against stoichiometric number, n, are usually irregular (Patrick et al., 1986; Yang et al., 1986a). Consequently, isopleths of log (σ/S cm−1) against 1000 K/T vary in gradient, and may cross each other. In some systems, e.g. PEOn:Ca(ClO4)2·6H2O, the conductivity maximum appears to lie at concentrations of salt that are more dilute (i.e. n > 18) than those studied. This inverse relationship between conductivity and salt content is also observed (Poulsen et al., 1985) for PEOn:LiClO4 and other polymer electrolytes. In Table 1.4, the compositions for maximum conductivity, nmax, are tabulated as a function of temperature for a range of systems. Inspection of Table 1.4 reveals that nmax varies substantially with the anion, cation and the temperature, from below 8 to above 16. The observed limits of the variation, i.e. 4–24, may be artificially narrow and merely reflect the truncated range chosen by individual authors. Values of nmin are presented in Table 1.5. Although also irregular, it can be seen that, except for some hydrated salt systems near room temperature, nmin < nmax. It is clear that nmax and nmin do not correlate with cationic radius (or charge density) although at certain temperatures it has been noted that the conductivity at fixed n decreases with increase in cationic radius (Patrick et al., 1986). As can be seen in Table 1.6, the conductivity results for PEO:X(ClO4)2·6H2O complexes, where X = Mg2+, Zn2+, Ca2+, Sr2+, are also in accordance with that observation. There appears to be some ambiguity regarding the choice of the anion that will yield maximum conductivity with a given cation, and vice versa. Further significant observations can be reported, however, as discussed below. Patrick and co-workers (1986) have found that the polyether electrolytes of divalent cation perchlorates (ClO4−) are generally better conductors than the corresponding thiocyanates (SCN−). Table 1.6 shows results
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Polymer electrolytes
Table 1.4 PEOn : salt ratios for maximum conductivities Temperature (° C) Salt
20
30
40
50
60
70
80
90
100
110
120
Ca(ClO4)2·6H2O Ca(SCN)2·3H2O Cu(ClO4)2 MgCl2 Mg(ClO4)2·6H2O Mg(SCN)2·4H2O PbBr2 Sr(ClO4)2·6H2O ZnCl2 Zn(ClO4)2·6H2O PbI2 PbCl2
12 9
12 9 8 16 12 18 8 12 8 18 24 16
18 12 8 16 12 18 8 12 8 18 24 16
18 12 8 16 12 18 8 12 8 18 24 16
18 15 8 16 12 18 8 12 8 18 24 16
18 15 8 16 15 18 8 12 8 18 24 16
18 15 8 16 18 12 8 12 4 18 24 16
18 15 8 16 18 12 8 12 4 18 24 16
18 15 8 16 18 12 8 12 4 18 24 16
18 15
18 15
16 15 12 8 12
15 15 12 8 12
18 24 16
18 24 16
12 15 12 18
Table 1.5 PEOn : salt ratios for minimum conductivities Temperature (° C) Salt
20
30
40
50
60
70
80
90
100
110
120
Ca(ClO4)2·6H2O Ca(SCN)2·3H2O Cu(ClO4)2 Mg(ClO4)2·6H2O Mg(SCN)2·4H2O Zn(ClO4)2·6H2O ZnCl2 MgCl2 PbBr2 PbI2
15 6
15 6 4 6 6 9 12 4 16 16
9 6 4 6 6 9 12 4 16 16
9 6 4 6 6 9 12 4 16 16
9 6 4 6 9 9 12 4 16 8
9 6 4 6 6 9 12 4 16 8
6 6 4 6 6 9 12 4 16 8
6 6 4 6 6 9 12 4 16 8
6 6 4 6 6 9 12 4 16 8
6 6 4 6 6 9 12 4 16 8
6 6 4 6 9 9 12 4 16 8
15 6 9
for the Mg and Ca systems confirming this observation, which seems to be also valid for other salts (e.g. the conductivity of PEO:Mg(ClO4)2·6H2O complexes is higher than the conductivity of PEO:MgCl2 complexes). The fact that PEO:Zn(ClO4)2.6H2O complexes (Patrick et al., 1986) present lower maximum conductivities than PEO:ZnCl2 complexes (Abrantes et al., 1986) may be due to the different preparation procedures of the polymers (different water content trapped in the polymer films), apart from the larger size of the ionic carriers. The hypothesis of ion-pairing between ClO4− and Zn2+ can also be advanced, but requires experimental confirmation.
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Introduction to polymer electrolyte materials
27
Table 1.6 Maximum conductivities of the PEO : salt complexes, given in (−log σ) values Temperature (° C) Salt
20
30
40
50
60
70
80
90
100
110
120
Ca(ClO4)2·6H2O Ca(SCN)2·3H2O Cu(ClO4)2 MgCl2 Mg(ClO4)2·6H2O Mg(SCN)2·4H2O PbBr2 Sr(ClO4)2·6H2O ZnCl2 Zn(ClO4)2·6H2O PbI2 PbCl2
8.0 8.8 4.7 9.4 6.4 9.2 9.0
7.2 8.2 4.3 8.8 5.6 8.3 8.6 7.5 4.7 6.7 9.6 9.1
6.2 7.5 4.0 8.0 5.0 7.0 8.2 6.8 4.4 5.8 9.2 8.7
5.2 6.6 3.4 7.2 4.7 6.2 7.8 6.3 4.1 4.9 8.4 8.2
4.8 5.8 3.2 6.5 4.3 5.9 7.4 5.7 4.0 4.6 7.4 7.6
4.7 5.5 3.0 5.7 4.1 5.6 7.0 5.3 4.0 4.3 6.5 6.7
4.4 5.3 3.0 5.5 4.0 5.4 6.0 5.0 3.7 4.0 6.4 6.5
4.1 5.2 2.9 5.2 3.9 5.1 5.9 4.7 3.6 3.8 6.3 6.4
4.0 5.0 2.8 5.0 3.8 5.0 5.8 4.4 3.4 3.8 6.2 6.3
3.9 4.9
3.8 4.8
4.9 3.7 4.8 5.7 4.2
4.7 3.6 4.7 5.5 3.9
3.7 6.1 6.3
3.6 5.9 6.2
5.3 7.6 9.4
As far as the effect of the different anions on the conductivity is concerned, it has been shown that (Abrantes et al., 1986; Huq et al., 1987; Yang et al., 1986a,b): σ(PEO:PbBr2) > σ(PEO:PbCl2) > σ(PEO:PbI2)
(T < 60 °C)
σ(PEO:PbBr2) = σ(PEO:PbI2) > σ(PEO:PbCl2)
(T = 60 °C)
σ(PEO:PbI2) > σ(PEO:PbCl2) > σ(PEO:PbBr2)
(T = 70 °C)
σ(PEO:PbBr2) > σ(PEO:PbI2) > σ(PEO:PbCl2)
(T > 70 °C)
Comparing these sequences with the values of the ionic radii for the anions Cl−, Br− and I−, it can be concluded that there is no relationship between maximum conductivities and the ionic radii. If we consider a fixed value of n, we can see that the increase in conductivity is a direct function of the increase of the cationic mobility. For example, the ionic radii for Sr2+, Ca2+, Zn2+ and Mg2+ decrease in the order Sr2+ > Ca2+ > Zn2+ > Mg2+ and the opposite should occur with the corresponding mobilities. Indeed, Table 1.5 shows that at 30 °C, for example, the maximum conductivity for the corresponding complexes increases in the order PEO12:Sr(ClO4)2·6H2O < PEO12:Zn(ClO4)2·6H2O < PEO12:Mg(ClO4)2·6H2O. An obvious exception to this rule is shown by complexes such as PEO:Cu(ClO4)2·6H2O and PEO:Mg(ClO4)2·6H2O. The ionic radius of Cu2+ is higher than that of Mg2+, so the σMg complex should be higher than the σCu complex. However, this is not the case, which can be easily understood
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Polymer electrolytes
on the basis of the fact that Mg2+ mobile ions are surrounded by their hydration shells. Moreover, the Mg2+ ions are apparently trapped in strong electrostatic bonds with the ether linkages on the polymer chains. Papke et al. (1982) have found that PEO–alkali metal salt complexes exhibit no detectable crystallisation for the larger cations such as Rb+ and Cs+, while Na+ salt complexes showed a high degree of crystallinity. A similar observation was noticed for PEO:metal salt complexes with larger anions. Furthermore, Watanabe and Ogata (1987) also reported that the larger the cation radius, the higher the conductivity values for PPO:SCN (alkali metals) electrolytes. Accepting that this phenomenon can also be applied to PEO:divalent metal salt, we can assume that large cations and/ or anions favour the production of complexes possessing large proportions of amorphous regions, which explains the higher conductivities obtained experimentally. In fact, the suppression of crystallisation improves markedly the transport properties even at room temperature where linear PEO is most crystalline, as we have shown (Araújo and Sequeira, 1993) by synthesising cell-amorphous interpenetrating polymer networks (Fig. 1.5). For the same complex, at a certain temperature, and variable n, we can observe that the conductivity generally increases as the salt concentration increases (n decreases), then it attains a maximum for intermediate n values, and finally it decreases for higher salt concentrations. This behaviour is typical of, for example, PEOn:Cu(ClO4)2, PEOn:ZnCl2 and PEOn:MgCl2 complexes. High salt concentrations have four main effects, they: (a) favour the formation of the PEO:salt complex, leaving a film depleted in free PEO (Yang et al., 1986b); (b) hinder the complete salt dissociation; (c) facilitate the
1 95 °C 85 °C 75 °C 65 °C 55 °C 45 °C 35 °C 25 °C
Log (σ /S cm–1)
0 –1 –2 –3 –4 –5 –6 0
8
16
24
32
40
n
1.5 Conductivity as a function of n in PEOn:Mg(ClO4)2. Reprinted with permission from Arau´jo and Sequeira, 1993. Copyright 1993, The Electrochemical Society.
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Introduction to polymer electrolyte materials
29
formation of ionic multiplets, such as triplets (X− M+ X−, M+ X− M+) and quartets (M+ X− M+ X−), which are less mobile than free ions and, in addition, decrease their effective concentration (Ratner, 1987); and (d) diminish the number of vacancies available for ionic transport (Sequeira, 1983a; Sequeira and Hooper, 1983b,c). These four effects explain the low conductivities observed at higher salt concentration. The low conductivities measured at low salt concentrations (n ≥ 12) are to be expected due to the reduced number of available charge carriers. Cation–anion interactions inside the polymer electrolyte, such as the phenomenon of ‘ion-pairing’ (Papke et al., 1981b), can explain conductivity variations over the whole range of temperatures and stoichiometries used. Salts whose anions are large, easily polarisable and weak Lewis bases, i.e. bases whose conjugate acids are strong, have a reduced tendency for ionpairing, and their incorporation in the polymeric material leads to higher conductivity complexes. Thus, the concentration dependence of the conductivity is a very complex function, as it is illustrated in Fig. 1.6, which relates an amorphous interpenetrating polymer network of PEO–LiClO4 (Hudson and Sequeira, 1993, 1995). A number of studies have been carried out where pressure rather than temperature has been used as a variable in conductivity measurements. Archer and Armstrong (1980) found that the bulk conductivity of a PEO4.5:LiCF3SO3 electrolyte decreased with increasing pressure over the range 60 to 300 MPa. When a cell was held at a constant pressure and
+
–4.0
+
Log σ (S cm–1)
+
+
+
–4.5 +
+
+
–5.0
+ –5.5
50
100 n
150
200
1.6 Variation in conductivity with salt concentration for an amorphous interpenetrating polymer network containing PEOn:LiClO4. Reprinted with permission from Hudson and Sequeira, J. Electrochem. Society 142(12), 4013–4017. Copyright 1995, The Electrochemical Society.
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Polymer electrolytes
temperature, the conductivity was again seen to fall with time. These effects were attributed to some form of structural change at higher pressures and may possibly involve increased crystallisation (Vincent, 1987a). Measurements of transport processes in crystalline solids as a function of pressure allow the activation volume, V*, for the conduction process to be evaluated. It is defined by (Flynn, 1972): V* =
( )
δG T δP
[1.6]
where δG represents the free energy change associated with the basic steps in the conduction mechanism. Values of V* can then be used to test theoretical models of the process. This approach has had application in a wide number of solid systems and has been applied to polymer electrolytes by Bridges and Chadwick (1988), Chadwick et al. (1983), Greenbaum et al. (1988) and Wintersgill et al. (1987), among others. Both isobaric and isothermal conductivity studies have been carried out. In determining the most appropriate method of association of the free energy change with the ionic conductivity, two options were discussed, one described by an Arrhenius expression, the other by a free volume expression. Interpreting trends in activation volumes with respect to charge transport is complex. In particular, it has been proposed that factors such as the possession of a permanent dipole by the anion (e.g. CF3SO3− or SCN−) may have a significant effect on its mobility (Lee and Crist, 1986) and thus in its contribution to conductivity.
1.4.3 Alternating current conductivity The problems of concentration polarisation which complicate DC measurements are largely avoided if AC is used instead. AC studies are similar to the DC technique in that the ratio of voltage to current is measured. For DC, this ratio provides the value of the resistance, R, measured in ohms (Ω). For AC the ratio gives an analogous quantity, the impedance, Z, also measured in Ω. The impedance contains four main contributions; these are from resistance, capacitance, constant phase elements (CPE) and inductance. The latter is unimportant for polymer electrolytes although it can play a role in other electrochemical applications of polymers. Measurement of the impedance as a function of frequency is called impedance spectroscopy. The detailed interpretation of this variation provides important information about the contributions to the test-cell electrochemical behaviour of its various components: interfaces, intra-electrolyte boundaries and the bulk electrolyte itself. This information is normally extracted by finding an idealised equivalent circuit, the calculated fre-
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Introduction to polymer electrolyte materials
31
quency-dependent behaviour which mimics the actual behaviour of the test-cell. The components (resistors, capacitors, etc.) of the equivalent circuit which are networked together in some combination of series and parallel connections can be identified with parts of the test-cell. The frequency dependence can be displayed in several ways, each of which emphasises a different frequency regime or a particular contribution. To appreciate these, it is necessary to recognise the fact which is not particularly palatable to the non-mathematically inclined reader, that impedance is a complex quantity. In practice, this means that it is made up of two parts: one called real, which is labelled Z′, and the other called imaginary and labelled as Z″. The contributions to Z″ are all terms which involve the quantity j, the square root of minus one (√−1). Z′ and Z″ are related to Z by: Z = Z′ − jZ″
[1.7]
The various resistors, capacitors, etc., which are connected together in series and parallel combinations to form the equivalent circuit that corresponds to the test-cell, contribute to the real and imaginary parts in a way that is quite easy to calculate. The frequency-dependent information can be displayed in the form of Z′ or Z″ as functions of the angular frequency, ω, or of log (ω). An alternative display, much favoured in the polymer electrolyte field, is the so-called complex plane representation, in which −Z″ (y-axis) is plotted against Z′ (x-axis). This display is characterised for polymer electrolytes by arcs, which have the form of depressed or flattened semicircles, and straight lines, inclined to the x-axis, which are usually called titled spikes. The advantages of this form of data presentation are that each arc or spike is characteristic of a particular region of the cell and that the bulk resistance can be easily read from the graph. A disadvantage is that the frequency of each measurement point is not immediately apparent. The simplest test-cell that can be envisaged consists of two blocking electrodes in contact with the polymer whose conductance is to be measured. The three-component equivalent circuit is shown in Fig. 1.4(a). The polymer acts as a resistor, Rb, which is in series with the double layer capacitor, Cdl, at the interface, and in parallel with the geometric capacitance, Cg, as discussed in Section 1.4.2. This simple network is dealt with by evaluating the parallel term first to give an impedance Zp, which is then added to the double-layer impedance. The RC parallel combination gives Zp =
R (1 − jωCg R ) 1 + (ωCg R )
2
and the total impedance Z is now
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[1.8]
32
Polymer electrolytes Z=−
R (1 − jωCg R ) j + 2 ωCdl 1 + (ωCg R )
[1.9]
so that Z′ =
R 2 1 + (ωCg R )
[1.10]
and ⎡1 + (ωCg R ) (1 + Cdl Cg )⎤⎦ Z ′′ = ⎣ ωCdl ⎡⎣1 + (ωCg R )2 ⎤⎦ 2
[1.11]
This is already a very complicated expression, although the equivalent circuit was the simplest that could be chosen to represent the physical situation of a conducting cell. The impedance plot is what might be intuitively expected from bringing together a series combination, which produces a vertical spike giving a real-axis value of R, and a parallel combination, giving a semicircle tending to the value of R at low frequency. The closer the values of Cdl and Cg (i.e. the thinner the electrolyte), the more the semicircle and the spike tend to merge, as shown in Fig. 1.7. If the electrodes are non-blocking, then the Cdl is now shunted in parallel by a charge-transfer resistance, as shown in Fig. 1.4(b) and discussed in Section 1.4.2. Evaluation of the impedance expression for this equivalent circuit produces two semicircles. The high-frequency semicircle is related to the bulk electrolyte and the low-frequency semicircle, which is more distant from the origin, arises from interfacial processes. The bulk resistance is the Z′ value at the high-frequency end of the interface semicircle.
–Z ˝
Cd/Cg = 25
Cd/Cg = 1000 Z´
1.7 The effect of the Cdl/Cg ratio on the impedance plot.
© Woodhead Publishing Limited, 2010
Introduction to polymer electrolyte materials 2.4 E4 2.4 E5
– Im Z (Ω cm2)
6.25 E3 – Im Z (Ω cm2)
2
– Im Z (Ω cm )
42 °C
0
70 °C
0
(a)
2.0 E5 Re Z (Ω cm2)
51 °C 0 (b)
128.5 °C
0
58 °C
24 °C 0
33
2.0 E4 Re Z (Ω cm2)
114 °C 0
5.00 E3 Re Z (Ω cm2)
(c)
1.8 Measured impedances for the cell Ni/PEO6:NiCl2/Ni at seven different temperatures (from Plancha et al., 1992a). Reprinted from Solid State Ionics, 58, Plancha MJC, Rangel CM and Sequeira CAC, AC conductivity of polymer complexes formed by poly(ethylene oxide) and nickel chloride, page 6, Copyright (1992), with permission from Elsevier.
Complex impedance plots for the Ni/PEO6NiCl2/Ni cell recorded at seven different temperatures are shown in Fig. 1.8. At temperatures up to 50 °C, the system exhibits two semicircles corresponding to the electrolyte impedance (parallel combination of the geometrical capacitance, Cg, and the electrolyte resistance, Rb) and to the interfacial impedance (charge transfer resistance, Rct, and double layer capacitance, Cdl, coupling). For temperatures above 50 °C, the high frequency semicircle does not appear; this is due to the fact that the corresponding characteristic frequency is higher than 65 kHz, the limit of the frequency range of the apparatus used. Rb and Rct values also decrease with increasing temperature. For temperatures from 70 °C up to 130 °C, the second semicircle, characteristic of intermediate frequencies, almost disappears and a diffusion process is evidenced by an arc with a 45° high-frequency branch, curving towards the real axis at lower frequencies (see Fig. 1.8c) (Plancha et al., 1992a). The value of ω at the top of the semicircle related to the bulk electrolyte is given by ω = 1/CgR
[1.12]
and increases substantially as the resistance falls (for example, with the temperature of the polymer sample), modifying the usual aspect of the impedance curves for polymer electrolytes. A feature of the impedance plot that is particularly pronounced for polymer electrolytes is the depression
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Polymer electrolytes
of the semicircle and the tilting of the spike. No combination of resistors and capacitors (or inductors) will produce either the depression or the tilting, but both of these are natural consequences of using CPE. CPE are hybrids between a resistor and a capacitor with an impedance of the form: Zcpe = K [cos(pπ/2) − j sin(pπ/2)]/ωp
[1.13]
where 0 ≤ p ≤ 1 and K is a constant. For a simple circuit consisting of a resistor and a CPE in series, the Z plot is represented by a line, inclined at an angle pπ/2 to the Z′ axis and contacting this axis when Z′ = R, as shown in Fig. 1.9(a). For a resistor and a CPE in parallel, a depressed semicircle is produced, as shown in Fig. 1.9(b). This cuts the Z′ axis at the origin and at Z′ = R. Because all angles subtended within a semicircle are right angles, the point Z′ = R lies vertically above the far end of the diameter of the semicircle. The length of this diameter, d, can be seen from Fig. 1.9(b) to be given by: d = R/sin(pπ/2)
[1.14]
–Z ˝
CPE pπ 2
R Z´
(R,O) (a) –Z ˝
(R,R) CPE
T X d 2
(R,O) Z´
C
R (b)
1.9 Impedance representation of the combination of a resistor and a constant phase element (a) in series and (b) in parallel.
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and the distance, CT, from the centre to the ‘top’ of the semicircle is d/2. The height of the top of the semicircle, XT, is clearly CT − XC, and since XC = [d cos(pπ/2)]/2
[1.15]
then XT =
d d R [1 − cos ( pπ 2 )] − cos ( pπ 2 ) = 2 2 2 sin ( pπ 2 )
[1.16]
The undepressed semicircle resulting from an RC combination in parallel obviously has a top at the point (R,R) and so the depression is given by R − XT.
1.4.4 Direct current electronic conductivity To determine the transport number of the electrons and electron holes in a polymer electrolyte, which may be a factor of 10−3 or less than that of the ions, the DC polarisation method of Wagner may be used. In this method, the electrolyte Pn:MX is sandwiched between an electrode of the parent metal, M, which acts as reversible electrode, and an ion-blocking electrode such as graphite or platinum. The ion-blocking electrode can exchange electrons with Pn:MX but not ions. A DC potential E, below the decomposition potential Ed, is applied to the symmetric cell. The positive pole is on the ion-blocking electrode (see Fig. 1.10). Because there is no source or sink of M or X at the ion blocking electrode, at steady state only electrons and holes migrate through the sample of Pn:MX. For an electrolyte in which tion = tcation = 1, the relevant equation is: I ss = I − + I + =
+
{
( )
( ) }
ART 0 ⎡ EF ⎤ EF + σ +0 ⎡exp σ − 1 − exp − − 1⎤ ⎢ ⎥ ⎢ ⎥⎦ ⎣ ⎣ LF RT ⎦ RT
Pn : MX
M C+ = C+0 C– = C0–
Graphite
[1.17]
–
C+ = C+0 exp (EF/RT) C– = C0– exp (–EF/RT)
1.10 Wagner’s DC polarisation cell. The negative electrode is a reversible electrode and the activity of M is unity at the M|Pn:MX interface. A DC potential E is applied so that the activity of M at the Pn:MX|ion-blocking electrode interface is exp(−EF/RT) at steady state. Boltzmann statistics for an ideal solid solution are assumed for the electrons and electron holes in Pn:MX. The concentration of electrons and electron holes are denoted by C− and C+, respectively. The superscript zero denotes that the concentration in Pn:MX is at equilibrium with M.
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Iss
0 0
E
1.11 Schematic diagram of Iss versus E for a polarised cell of the type shown in Fig. 1.10.
where Iss is the steady-state electronic current, I− and I+ are the current due to electrons and holes, respectively, F is the Faraday’s constant, σ0− and σ0+ are the partial conductivities due to electrons and holes, respectively (the superior zero denotes equilibrium with pure metal, M), E is the applied voltage, A is the cross-section area of the sample and L is its thickness. A schematic diagram of the Iss–E curve for such a polarised cell is shown in Fig. 1.11. Generally, the equation for Iss may be simplified, yielding
( )
I ss ART ⎡ 0 EF = σ + exp + σ −0 ⎤ ⎦⎥ 1 − exp ( − EF RT ) LF ⎣⎢ RT
[1.18]
from which a plot of Iss/[1 − exp(−EF/RT)] versus exp(EF/RT) yields, in principle, the values of σ0− and σ0+. However, within each volume element there is equilibrium between the electrons and electron holes; that is, the intrinsic constant, k, is satisfied. Hence, the product C− C+ is very small and σ0−, σ0+ and exp(−EF/RT) are small. Accordingly, the plot from equation [1.18] rarely yields a positive slope and positive intercept within the experimental uncertainty of the data. The use of AC measurements to obtain the total conductivity and the DC polarisation cell to obtain the electronic conductivity allows the separation of ionic and electronic conductivity whose values may differ in many orders of magnitude. A key feature of the Wagner cell methodology is that the activity of the mobile ionic species is controlled at the reversible electrode, and for many polymer electrolytes both anions and cations conduct (Abe et al., 2004; Chandrasekhar, 1998; Hu et al., 2007; Itoh et al., 2009a; Judeinstein et al., 2008; Martins and Sequeira, 1990; Plancha et al., 1992a; Rajendran et al., 2008; Saito et al., 2003; Singh and Bhat, 2004; Stephan, 2006; Xi and Tang, 2004; B. Zhang et al., 2006). It is, therefore, only for polymer electrolytes in which ionic transport is effectively entirely either anionic or cationic that the use of the Wagner technique can be usefully contemplated. An alternative to the Wagner
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polarisation technique is the use of transient methods. Electrons and holes respond much more rapidly to fluctuations in an applied potential than ions do. A cell of the type M|electrolyte|M has a DC voltage applied across it and the current–time characteristics are monitored as a rapid voltage pulse is superimposed. The response will be rapid for a sample with a high electronic conductivity and slow for a good electronic insulator.
1.4.5 Ionic conductivity There are many AC and DC techniques for probing the transport of charged and neutral species. Among them, we can refer to the electrolysis method, the concentration cell techniques, and the ionic polarisation cell technique. The electrolysis method consists of studying the changes in the amount of material in the anode and cathode compartments of a cell. For liquids, this is done by the Hittorf method in which concentration changes are measured by titration, subsequent to the passage of a known total amount of charge. The technique has been applied to low molar mass polymer electrolytes by Cameron et al. (1989), among others. For solids, this type of technique is known as the Tubandt method, and involves the careful weighing of easily separated, non-adhering discs of electrodes and electrolyte material (Linford and Hackwood, 1981). This type of polymer material for which adhesion between the cell layers has been successfully avoided belongs to the group of network crosslinked materials studied by Cheradame and Niddam-Mercier (1989). In both the Hittorf and Tubandt methods, it is necessary to verify that the electrolyte compartment remains unchanged. The concentration cell techniques involve the use of concentration cells. Bouridah et al. (1986) measured the electromotive force (EMF) of suitable cells under conditions where no current was passed, and obtained the anion transference number from τ− =
−FdEcell RTd ln a
[1.19]
where a is the activity of the salt. Such measurements require an independent determination of the variation of the activity with the concentration. To avoid this, later studies on the voltages of similar cells with and without transference were carried out (Bouridah et al., 1988). The ionic polarisation cell technique involves transport under a chemical potential and electrical gradient and, therefore, includes the contributions of both charged and neutral species. The AC procedure has been suggested by Sorensen and Jacobsen (1982) and is based on the theory of MacDonald (1973, 1974) which involves analysis of the impedance spectra of the symmetric cell M|Pn:MX|M.
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Rb
Re
Z ˝ (ohms)
Zd
Cb
Ce
f
Rb
Rb+Re
DC limit
Z´ (ohms)
1.12 Idealised equivalent circuit and complex plane plot for a polymer electrolyte cell with electrodes that are non-blocking to the cation. Re = electrode resistance; Ce = electrode capacitance; Rb = bulk (or electrolyte) resistance; Cb = bulk (or electrolyte) capacitance; Zd = diffusion-controlled impedance.
For the idealised case, the equivalent circuit and response of Fig. 1.12 may be anticipated. The high-frequency semicircle is due to cation migration through the bulk and dielectric polarisations. At lower frequencies, a semicircle arises as a result of the charging and discharging of the electrode– electrolyte interface and reactions of the electroactive ion at the electrode interface. At the lowest frequencies, the current is affected by concentration gradients that give rise to diffusion in the electrolyte, and a skewed semicircle that approaches the limiting DC value of the cell impedance is observed. For fully dissociated electrolytes, the cation transference number τ+ can be evaluated by comparison of the width of the skewed semicircle, the diffusion-controlled impedance, Zd, with the value of the bulk resistance, Rb: τ+ = 1/(1 + Zd/Rb)
[1.20]
The AC technique has been by far the most widely used method for obtaining transference numbers on polymer electrolytes (Plancha et al., 1991, 1992b,c, 1993a,b, 1994, 1997c, 1998, 1999; Sorensen and Jacobsen, 1982; Yang et al., 1986a,b), but again the analysis is relevant only when mobile associated species are not present. In addition, electrode phenomena such as finite electrode kinetics or formation of passivating layers are not con-
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sidered and, as suggested by Fauteux (1985), the diffusion responsible for the low frequency arcs may not be a bulk process but one of the passivating films on the electrode surface, which are generally observed when polymer electrolytes are in contact with metallic lithium. The preceding considerations have dealt with simpler situations. The effects of non-ideality and higher associated charged clusters have yet to be rigorously analysed. The extent to which the preceding techniques will lead to results deviating from the real situation remains unknown.
1.5
Applications in practical devices
The future prospects for polymer electrolytes look promising because it has been appreciated that they form an ideal medium for a wide range of electrochemical processes. Other than primary and secondary batteries, and high and low temperature fuel cells, practical applications for polymer electrolytes that are under consideration include electrochromic devices, modified electrode/sensors, solid-state reference electrode systems, supercapacitors, thermoelectric generators, high-vacuum electrochemistry and electrochemical switching. Device applications that have been the main driving force behind the development of polymer electrolytes are treated hereafter.
1.5.1 Batteries Recently there has been a phenomenal growth in electronic devices and they are all demanding in terms of electricity supply. The only way of achieving this in practice in the foreseeable future is by a rechargeable battery. Batteries will play a key role in the development of new, self-powered medical implants, in ‘hybrid’ vehicles that run alternately on combustion engines and batteries, and in energy storage for renewable power sources such as wind and solar. Each type of battery will require different characteristics in terms of the storage capacity, the number of times the battery can be discharged and recharged, and the cost of the materials. The concept of the rechargeable battery is elegantly simple. Two electrodes are separated by an electrolyte that is an ionic conductor and electronic insulator. Charging the battery involves applying an electrical potential to force cations – lithium, say – up an energy gradient into the anode, where the ions are intercalated within a crystalline lattice. When the two electrodes are connected via an external circuit this pent-up chemical energy is released: the positive ions drift across the electrolyte and fall into an energy well within the cathode; at the same time, electrons flow to the cathode via the external circuit, providing an
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electric current. Within this system there are several ways to improve the battery, depending upon the priority required for the given application. Most water-based battery systems, such as nickel–cadmium, where the electrolyte is in an aqueous solvent, are limited to around 1.5 to 2 volts, whereas a lithium battery can give 3 to 4 volts per cell. Since lithium batteries can store more energy than other systems, most research effort on rechargeable batteries is focused on lithium technology (Koksbang et al., 1994; Sequeira, 1983b, 1987; Sequeira and Hooper, 1985; Sequeira and Marquis, 1986; Stephan, 2006). Rechargeable lithium ion batteries were introduced commercially by the Sony Corporation in the mid-1990s. A lithium cell that sits in a laptop computer typically consists of a graphite anode and a cathode whose ‘active’ ingredient is a lithium transition metal oxide, such as LiCoO2, with the two electrodes separated by an electrolyte of LiPF6 dissolved in an organic solvent, or an SPE (Aravindan et al., 2008; Itoh et al., 2009b; Jeon and Kwak, 2006; Kang et al., 2009; Koksbang et al., 1994; Meyer, 1998; Reiter et al., 2009a; Seki et al., 2005; Stephan and Nahm, 2006; T. Zhang et al., 2008). There are two ways to increase the amount of energy that such a cell can store. One is to increase the lithium mobility in the system – the more ions that can make their way from the anode to the cathode, the greater the storage capacity of the battery. This means materials that can store more lithium per unit weight and volume. The other approach is to increase the cell voltage. So, materials chemists are looking for new electrode materials that can store more lithium at higher voltage in a way that is inexpensive, safe and not toxic. The system must also be capable of being cycled – charged and discharged – hundreds or thousands of times with no appreciable loss of efficiency. Since the introduction of the lithium cell, most improvements to the technology have relied on engineering rather than chemistry. For example, the positive electrode consists not only of the lithium transition metal oxide, but also of conducting additives and binders. Over the years, new manufacturing processes have allowed the proportion of the inactive components of the electrodes to be lowered, giving way to a greater quantity of the lithium-containing compound. But this has effectively reached its limit and new approaches are needed. In current lithium batteries based on LiCoO2 cathodes, only ‘half’ a lithium ion is available for shuttling between the electrodes. Most recent work worldwide has focused on how to get that extra half lithium into the system. However, this will only double the charge stored and more radical approaches are required to go further. One approach is to remove the intercalation host altogether, and let the lithium react directly with oxygen in the air. The lithium entering the porous electrode reacts directly with gaseous oxygen in the air contained within the pores to form lithium oxide.
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Early results have shown that such an oxygen electrode can store significantly more charge than conventional systems relying on intercalation chemistry. Storage is expressed in milliampere hours per gram (mA h g−1), and LiCoO2 can typically store 140 mA h g−1, with a theoretical maximum of some 300 mA h g−1. Greater storages can be obtained by combining chemistry and novel manufacturing techniques. For example, researchers at St. Andrews (Scotland) have synthesised LiNi1/3Co1/3Mn1/3O2 using a new route that results in a macroporous material delivering a higher capacity at a higher rate than is currently possible (Shaju and Bruce, 2006). Given that the ions shuttle between one electrode and another, the ions that emanate from the cathode must have a home to go to in the anode. In other words, the anode’s storage capacity must be able to match that of the cathode. So far this has not been a problem: graphite can store around 300 mA h g−1. However, if the cathode’s storage capacity starts to outstrip this, the imbalance will need to be addressed. Sony Corporation has already improved this aspect of the negative electrode by using an alloy of lithium and tin, which can store between 700 and 1000 mA h g−1. Along with the storage capacity, a key characteristic of a rechargeable Li ion battery is just how rechargeable it is, i.e. how many times it can be cycled without appreciable loss of capacity. This is intimately connected with phenomena at the interfaces between the electrodes and the electrolyte (North et al., 1982; Sequeira, 1983c, 1985; Sequeira and Hooper, 1983c,d,e,f; Sequeira et al., 1984). A typical battery system operates by transferring lithium ions from a transition metal oxide host into graphite. During the lifetime of the battery, certain degradation products are collected at the electrode surfaces. These can lead to unwanted reactions, in which the heat generated can in turn cause further degradation, resulting in a process called ‘thermal runaway’. This can bring about potentially hazardous instabilities within the system. Paradoxically, the initial stage of the degradation of the electrolyte is critical in stabilising the electrodes, but excessive subsequent degradation is undesirable. At the surface of the electrodes, lithium ions react with the electrolyte to form a layer of typically lithium fluoride, lithium carbonate or organic lithium compounds. If this layer, called the passivation layer, did not form, the graphite negative electrode would delaminate and eventually disintegrate. The important thing is that this layer must remain permeable to lithium ions during subsequent cycling of the battery. It is a bit like a brick wall with some bricks missing – a supporting structure that still allows lithium to go through. However, the presence of the layer increases the resistance of the battery, which in turn generates heat – something that must be minimised. One way of mitigating thermal runaway is to use additives in the electrolyte so that the initial passivation layer is more stable and less likely to grow upon subsequent cycling. Such additives can include large
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organic molecules that degrade to form a stable passivation layer; alternatively, complexing agents can be placed at the interface that chemically scavenges ions or molecules within the electrolyte to form a stable film. It is important to minimise the loss of capacity during the lifetime of the battery by improving the efficiency of ion transport in and out of the electrodes. If lithium ions react with products within the electrolyte to cause this layer to grow, some lithium is effectively removed from the process. Ideally, we want the same amount of lithium to move in and out of the electrodes on the thousandth cycle as on the first cycle. If we lose only 1% lithium on each cycle, we would not have much capacity left after a few hundred cycles. Apart from the lithium polymer electrolyte batteries as a whole, other battery technologies have been studied, as reported by several researchers (Abrantes et al., 1985; Araújo et al., 1994a,b; Bruce, 1995a; Gray, 1997; Patrick et al., 1987). Polymer electrolyte batteries are presently produced at a rate of several million units per month and are rapidly replacing the bulkier and less energetic nickel–cadmium and nickel–metal hydride batteries in popular devices, such as cellular phones and computers. In addition, lithium ion batteries are also scaled-up in view of their use in electric vehicles. However, although a commercial reality, polymer electrolyte batteries are still the object of intense research with the aim of further improving their properties and characteristics (Araújo and Sequeira, 1991, 1993, 1995a,b; Araújo et al., 1996b; Bruce, 2005).
1.5.2 Fuel cells Fuel cells constitute an attractive class of renewable and sustainable energy sources that are alternatives to conventional energy sources, such as petroleum, which have finite reserves. Energy generation from petroleum oil and natural gas through combustion in a heat engine being subjected to the Carnot cycle limitation is inherently inefficient and accompanied by environmental pollution. In contrast, a fuel cell is intrinsically energy efficient, non-polluting, silent and reliable. In some embodiments, it is a low temperature device that provides electricity instantly upon demand, and exhibits a long operating life. Energy efficiencies of about 50–70% can be achieved with fuel cells. Fuel cells combine the advantages of both combustion engines and batteries, at the same time eliminating the major drawbacks of both. Similar to a battery, a fuel cell is an electrochemical energy device that converts chemical energy into electricity and, akin to a heat engine, a fuel cell supplies electricity as long as fuel and oxidant are supplied to it. Among the various types of fuel cells developed so far, polymer electrolyte fuel cells (PEFCs) have the advantage of high power densities at relatively low operating temperatures (≤80 °C) (Antolini, 2004) and therefore
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are considered promising power sources for portable and residential applications. Research and development on PEFCs using hydrogen as the fuel have progressed enormously but their successful commercialisation is restricted due to problems with the safety and storage efficiency of the flammable hydrogen gas (Igarashi et al., 1995; Shukla et al., 2003). In order to overcome these difficulties, liquid methanol can be used instead to fuel PEFCs. Direct use of liquid fuel in a PEFC simplifies the engineering issues, thereby driving down the system complexity and hence the cost. PEFCs that are fed with methanol as the fuel are referred to as direct methanol fuel cells (DMFCs). However, DMFCs possess several limitations in terms of the high cost of the noble-metal electrocatalysts, low open circuit voltage (OCV), and methanol crossover from anode to cathode compartments (Shukla et al., 2003, 2004). Moreover, CO is formed during the inefficient methanol electrooxidation, which leads to poisoning of the platinum anode (Ganesan and Lee, 2005). The problems arising from the use of methanol in DMFCs can be overcome by using other hydrogen carrying materials as the fuel, such as various borohydride (BH4−) compounds. Sodium borohydride (NaBH4), which has a charge capacity value of 5.67 A h g−1 and a hydrogen content of 10.6 wt% (Lee et al., 2006; Santos, 2006, 2009; Santos and Sequeira, 2007) is a good alternative to methanol as a fuel. A PEFC that uses a BH4− compound, usually NaBH4 in aqueous alkaline medium, directly as a fuel is termed as direct borohydride fuel cell (DBFC). Although the concept of the DBFC was first demonstrated by Indig and Snyder (1962) in the early 1960s, Amendola et al. (1999) were the first to report a direct borohydride–air fuel cell that employed an anion exchange membrane (AEM) as electrolyte and exhibited a maximum power density of 60 mW cm−2 at 70 °C. The DBFC reported by Amendola et al. (1999) could use about seven out of a theoretically maximum eight electrons (per BH4− ion), exhibiting high fuel utilisation efficiency. DBFCs are considered attractive energy suppliers, especially for portable applications. DBFC supersedes DMFC in terms of capacity value, electrochemical activity, theoretical OCV (1.64 V for DBFC employing oxygen as oxidant as against 1.21 V for DMFC using oxygen as oxidant) and power performance at room temperature. In addition, use of alkaline electrolytes which feature relatively low corrosion activity, opens up the possibility of applying readily available and lowcost non-precious metal anode catalysts (Santos and Sequeira, 2010; Santos et al., 2007). Both BH4− and its oxidation product, metaborate (BO2−), are relatively inert and non-toxic. BO2− can be recycled to produce BH4− and the techniques involved are under investigation (Fakiogˇlu et al., 2004; Sequeira et al., 2007). One issue that needs further investigation is the development of polymer electrolyte membranes whose structure–property relationships increase their applicability in fuel cell systems, including PEFCs, DMFCs and DBFCs.
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The evolution of membranes for fuel cell applications started as early as 1959, by General Electric (GE), with the testing of phenolic membranes prepared by polymerisation of phenol-sulphonic acid with formaldehyde. These membranes had low mechanical strength and a short lifetime of 300–1000 hours and showed a power density of 5–10 mW cm−2 (Magnet, 1968). During 1962–1965, GE attempted to improve the power density by developing partially sulphonated polystyrene sulphonic acid membranes (prepared by dissolving polystyrene sulphonic acid in ethanol-stabilised chloroform followed by sulphonation at room temperature). This membrane exhibited a better water uptake and an improved power density of 40–60 mW cm−2, which enabled its application in NASA’s Gemini flights (Bockris and Srinivasan, 1969). Although initial attempts were unfruitful, GE subsequently redesigned their PEFC, and the new model P3, despite malfunctions and poor performance on Gemini 5, served adequately for the subsequent Gemini flights. However, this membrane exhibited brittleness in the dry state. Another approach to improve the mechanical strength and the membrane life was again undertaken by GE in the late 1960s by preparing crosslinked polystyrene–divinylbenzene sulphonic acid membrane/ polymer in an inert matrix. The life of the membrane ranged from 103 to 104 hours and the power density attained 75–80 mW cm−2 (Grott, 1985). The main problem encountered with all the above-mentioned types of membranes was that the proton conductivities were not sufficiently high to reach a power density even as low as 100 mW cm−2 (Wakizoe et al., 1995). In the 1970s, DuPont developed a polymer based on PFSA called Nafion that not only showed a two-fold increase in the specific conductivity of the membrane but also extended its lifetime (104–105 hours). This soon became a standard for PEFC and remains so today. The Dow Chemical Company and Asahi Chemical Company synthesised advanced PFSA membranes with shorter side chains and a higher ratio of SO3H to CF2 groups (Costamagna and Srinivasan, 2001). Table 1.7 (adapted from Smitha et al., 2005) provides a comparison of some commercial cation exchange membranes (CEMs). Many more proton polymer electrolytes have been developed for PEFCs and DMFCs as described in Section 1.2.2 and reported by Chen and Sequeira (2002), Gottesfeld and Fuller (1999), Hartnig et al. (2008), Lee et al. (2006), Neburchilov et al. (2007), etc. Concerning DBFCs, both the AEM and the CEM have served in a majority of fuel cell operating conditions (Atwan et al., 2005; Choudhury et al., 2005; Liu et al., 2003, 2005; Ponce de Leon et al., 2007; Santos and Sequeira, 2009a; Sequeira and Santos, 2009). In particular, Nafion 117 relying on the sodium ion for the ionic conductivity, instead of the hydrogen ion as in the case of the hydrogen PEFC, is often employed (Santos and Sequeira, 2006, 2009b). But development of low-cost polymer membranes with little BH4− crossover continues to be a significant key in bringing DBFC to the level of common usage.
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Table 1.7 Properties of commercial cation exchange (adapted from Smitha et al., 2005). Reprinted from J Membrane Sci, 259, Smitha B, Sridhar S and Khan AA, Solid polymer electrolyte membranes for fuel cell applications – a review, page 13, Copyright (2005), with permission from Elsevier
Membrane designation
Membrane type
Conductivity (S cm−1) Thickness Gel water at 30 °C and IEC (%) (mequiv g−1) (mm) 100% RH
Asahi Chemical Industry Company Ltd (Japan) K 101 Sulphonated 1.4 0.24 polyarylene
24
0.0114
0.15
25
0.0051
0.15
–
0.0071
0.15
–
–
Ionac Chemical Company, Sybron Corporation (USA) MC 3470 – 1.5 0.6 MC 3412 – 1.1 0.8
35 –
0.0075 0.0114
Ionics Inc. (USA) 61AZL386 – 61AZL389 – 61CZL386 –
2.3 2.6 2.7
0.5 1.2 0.6
46 48 40
0.0081 – 0.0067
DuPont Company (USA) N 117 Perfluorinated N 901 Perfluorinated
0.9 1.1
0.2 0.4
16 5
0.0133 0.01053
Pall RAI Inc. (USA) R-1010 Perfluorinated
1.2
0.1
20
0.0333
Asahi Glass Company Ltd (Japan) CMV Sulphonated 2.4 polyarylene DMV Sulphonated – polyarylene Flemion Perfluorinated –
Note: IEC stands for the membranes’ ion-exchange capacity and RH stands for the relative humidity.
1.5.3 Photoelectrochemical devices A photoelectric energy converter is based on semiconductors as active components. These may be crystalline, polycrystalline or amorphous. They may be inorganic, such as the well-established silicon semiconductor, or organic. Absorption of photons of requisite energy – those with energy exceeding the bandgap of the material – leads to the generation of electron–hole pairs, which can be separated by diffusion or drift, in a sufficiently strong electric field. The electric field is due to the double layer found at the junction of two dissimilar materials. These materials may be the same kind but with different doping, as found for instance in the p–n junction
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cells constructed from single crystal silicon. They may also be a semiconductor in contact with a metal, as in the Schottky barrier cell. The third class of cells is constructed from the junction of a semiconductor with an electrolyte. These are called photoelectrochemical cells (PECs) (Nozik, 1978). In a PEC, light is incident on the n-type semiconductor/electrolyte junction (photoanode), where light absorption occurs and an electron–hole pair is formed. The pair is separated by the strong electric field found just beneath the semiconductor surface, and the hole is driven towards the interface between semiconductor and electrolyte. Charge transfer to the redox species A contained in the electrolyte results in the oxidation to A+. Conversely, the electron is driven to the metal/electrolyte interface (counter electrode), where the redox species is reduced. No net chemical work is done and we can extract the energy as a current from the cell. The potential advantages of a photoelectrochemical energy converter are: • •
•
the ease of formation; the barrier is formed by contacting the semiconductor with the electrolyte; the possibility of forming high barriers by using a redox couple of extreme potential; these are reflected in high open circuit photovoltages; the possibility of using lower semiconductor materials; polycrystalline materials are easier to use than in solid state photoelectric devices (Hodes et al., 1984).
The first use of SPE containing redox species for use in PECs was in a tandem cell (Skotheim, 1981) and as contacts to silicon (Inganas et al., 1986; Skotheim and Inganas, 1985; Skotheim and Lundstrom, 1982). Other materials have also been used in combination with polymer electrolytes (Cook and Sammels, 1985; Sammels and Ang, 1984; Sammels and Schmidt, 1985). The requirements for application of polymer electrolytes in PECs is, of course, that they can complex desired redox species, that their conductance is high enough (due to high conductivity and/or small thickness), and that their mechanical and electronic properties are such that interfaces with desired electronic properties can be formed between the electrolyte and the semiconductor. We also require the optical absorption to be sufficiently low within the polymer film, so that negligible conversion losses are introduced in this manner. They should furthermore be stable over considerable periods of time, and should not allow deleterious reactions at the interface between the polymer film and the semiconductor. Practical aspects of manufacture and encapsulation strongly recommend the solid state approach to PECs. Polymer electrolytes are easily processed into thin films that can be spun or cast over large areas, which are some requirements for photoelectrochemical materials. They are encapsulated
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more easily than PECs using liquids, as the electrolyte does not seep out. They might have to be protected from the atmosphere, but this has been handled. Much work is needed for the development of rugged interfaces between the various phases in a PEC, including those between the polymer electrolyte and the semiconductor as well as those between the polymer electrolyte and surface bound electrocatalysts (Skotheim, 1986a). Nafion, PEO, and many other polymer-based electrolytes have been widely used in PECs, which show better performances than their counterparts using liquid electrolytes (Li et al., 2006; Skotheim, 1986b; Wang, 2009).
1.5.4 Electrochemical capacitors In recent years there has been increasing interest in the power capacitors, ultracapacitors or supercapacitors based on electrochemical systems. These include electric double layer capacitor (EDLC) types based on carbon electrodes with suitable electrolyte systems, and electrochemical capacitors with pseudocapacitance (Conway, 1991, 1995). Several types of capacitors are classified in Table 1.8 according to their energy storage mechanisms.
Table 1.8 Classification of capacitors according to the energy storage mechanism Basis of charge or energy storage
Type
Regular dielectric capacitor Electrostatic Vacuum Dielectric Electrolytic capacitor Electrostatic Oxide-film EDLC Electrostatic Aqueous electrolyte Non-aqueous electrolyte Ion-conducting polymer Solid inorganic salt Redox capacitor Redox oxide film Redox polymer film Soluble redox system Underpotential deposition
Example
– – Al, Ta oxide condenser
Carbon, activated carbon electrode, aqueous H2SO4 (Charge separation TEABF4/propylenecarbonate at double layer PAN gel electrolyte electrode interface) RbAg4I5 Faradic charge transfer RuO2·xH2O, IrO2, NiO/Ni, TiS2/Li+ (Pseudocapacitance) Polypyrrole, polythiophene, polyacene [Fe(CN)6]3−/[Fe(CN)6]4− H/Pt, Pb/Au (adsorption)
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Redox capacitors function by using electrochemical pseudocapacitance based on fast Faradic redox (Conway, 1991), insertion of protons into noble metal oxides (Liu and Anderson, 1996; Zheng et al., 1995), p- and n-doping of conducting polymers (Arbizzani et al., 1995), lithium ion intercalation (Passerini et al., 1995) or underpotential deposition (adsorption) processes (Conway and Tilak, 1992). Among the various materials for redox capacitors, proton insertion into noble metal oxides and p- and n-doping of conducting polymers have been widely investigated. Conway (1991) reported the use of RuO2 and IrO2 as electrodes of redox capacitors. Zheng et al. (1995) found the specific capacitance, 760 F g−1, of amorphous hydrous ruthenium oxide, where the process can be expressed as RuO2·xH2O + γH+ + γe− ↔ RuO2−x−γ(OH)2x+γ
0≤γ≤2
[1.21]
Noble metal oxides, however, are expensive and usually have poor electrical conductivity. The process of charge storage in conducting polymers requires electronic transport through the polymer backbone. Charging commonly used conducting polymers corresponds to one charge unit per 2–3 monomer units at the maximum. Conducting polymers belong to a class of materials with exciting potential in the field of supercapacitors because they can be charged and discharged at high rates. The principle of EDLC operation is very simple and is based on the wellknown electrical or double layer phenomenon. The device operates within a potential range in which no Faradic reactions take place, and thus the behaviour is fully capacitive. Polarisation of the electrodes in opposite directions leads to accumulation of opposite charges at the electrode–solution interfaces. The higher the electrode surface area and the polarity of the electrolyte solution and its ionic concentration, the higher the capacity and energy density of these devices. The capacitance (C) and the accumulated electrostatic energy (E) stored are given by equations [1.22] and [1.23], respectively (Nishino et al., 1985): C=∫ E=
ε dS 4πδ
[1.22]
1 CV 2 2
[1.23]
EDLC with very low internal resistance (no electrochemical reaction resistance), very high capacitance and extremely long cycle life (ca. 106 cycles) can be achieved (Murphy and Kramer, 1994). Thus, the EDLCs have been envisaged as pulse-power sources, an especially useful complement to batteries for electric vehicles. Improving the energy and power density of EDLC requires development of new electrode materials and electrolyte systems as well as new cell
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design. It appears from our own experience that using thin films of SPE in EDLC may be useful in achieving the advanced goals for ultracapacitors. Ion conducting polymers may be preferable in these devices’ electrolytes because of their flexibility, malleability, easy manufacture and chemical stability (for the same reasons that they have been applied to lithium secondary batteries). The gel electrolyte systems, which consist of a polymeric matrix, organic solvent (plasticiser) and supporting electrolyte, show high ionic conductivity of about 10−3 S cm−1 at room temperature and have sufficient mechanical strength (Reiche et al., 1995). Therefore, the gel electrolyte systems are superior to SPEs and organic solvent-based electrolytes as materials for batteries and capacitors, for room temperature operation (Choudhury et al., 2009; Kalpana et al., 2006; Kumar and Bhat, 2009; Zhang et al., 2009). The application of SPE to EDLC is attractive, but there have been few attempts in this direction for the following reasons: (1) relatively low ionic conductivity of the most relevant polymer electrolyte systems at room temperature; (2) poor contact at the electrode/electrolyte interface; and (3) detrimental influence of crystalline domains of the polymer electrolyte systems on their conductivity. Thus, many researchers focus on improving the ionic conductivity and increasing the electrode’s capacitance by controlling the electrode/electrolyte interface. The proton conducting polymer (Nylon 6-10, 2H3PO4) has been used as an electrolyte in a power capacitor (Lassagues et al., 1995), and this polymer blend associated with activated carbon gives an electrostatic capacity similar to that obtained when using liquid electrolyte solutions. However, the energy density of a capacitor with proton conducting polymer is limited to a voltage of about 1 V (which is the maximum voltage applicable for such systems). Non-proton-conducting SPEs such as PEO or PPO salts have a low dielectric constant (e.g. ε = 5 for PEO) and low ionic conductivity at room temperature (<10−5 S cm−1) (Araújo and Sequeira, 1995a; Sequeira and Marquis, 1988; Sequeira et al., 1996). However, such non-proton-conducting polymers have the advantage of presenting a large voltage region. Also, they can be manufactured as ultrathin films and thus high energy density may be obtained in devices containing them. Some time ago, the application of polymer electrolytes such as PEO–LiClO4 SPE for electrochemical devices was attempted (Gray, 1997; MacCallum and Vincent, 1989; Sequeira et al., 1994a). The phase diagram and ionic conductivity of PEO–LiClO4 electrolyte systems have been measured by many researchers (Hudson and Sequeira, 1995; Robitaille and Fauteux, 1986). This polymeric electrolyte consists of crystalline complexes and an amorphous phase. In fact, only the amorphous phase is conducting (Sequeira et al., 1993). However, there are few reports on the use of these electrolyte systems for EDLC. Osaka and co-workers have attempted using PEO–LiClO4 SPE as an electrolyte medium for EDLC.
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After a survey of carbon materials as electrodes for these devices, it was found that isotropic high density graphite (HDG) is a good candidate for the electrode material for EDLC of high capacity using a solid-state electrolyte (Liu and Osaka, 1996; Liu et al., 1996). The effects of some properties of EDLC composed of PEO–LiClO4 SPE and HDG electrodes were investigated. Preliminary studies performed recently at our laboratory on all-solidstate EDLC with PEO–LiClO4 SPE (Li : O = 1 : 8) at 25 °C showed a good charge/discharge behaviour over 1500 cycles between 0.0 and 3.0 V. In addition, various gel electrolytes were investigated for EDLC as well. Poly(ethylene oxide) (PEO)-, polyacrylonitrile (PAN)-, polymethylmethacrylate (PMMA)- and poly(vinylidene fluoride) (PVdF)-based gel electrolytes, which have been applied for secondary lithium batteries (Liu and Osaka, 1997; Schmutz et al., 1994; Sequeira et al., 1994b), were also studied as electrolyte systems for EDLC with carbon electrodes.
1.5.5 Electrochromic devices Electrochromic materials have been extensively studied for potential use in an information display and they are currently being used in automotive applications (glare-free automotive rear view mirrors, sunroofs) and find future applications for solar control (smart windows, etc.). Electrochromism (a solid-state effect) and electrochemichromism (an electrochemical effect) are exhibited by a number of materials, both inorganic and organic liquids and solids. Many of the systems described herein are based on electrochemichromism; however, no distinction will be made hereafter. Electrochromism is a reversible colour change in a material caused by an applied electric field or current. Ion insertion materials, which are mixed conductors (i.e. both electronic and ionic), into which ions can be rapidly and reversibly inserted, can be used as electrochromic materials. The optical properties of these materials depend on the number of inserted ions. They may include organic (diphthalocyanine and tetrathiafulvalene) as well as inorganic (WO3, Ir2O3, Ni(OH)2) materials. Among these, electrically conducting organic polymers are attracting considerable attention mainly because of their various electronic properties (Sequeira et al., 1995). In addition to their use as electrochromic conducting polymers, advanced polymer materials can be used as thin film electrolytes in electrochromic devices for controlling the reversible optical absorption associated with the applied potential (Sequeira and Santos, 2009). Both kinds of advanced materials – electrochromic conducting polymers and SPEs – have been intensively investigated in recent years. These materials can be used in several major electrochromic applications, namely display devices, reflective mirrors and ‘smart’ windows, since their optical
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absorption occurs in both the visible and near-infrared regions of the spectrum (Bozkurt, 2002; Marcilla et al., 2006; Nguyen et al., 2009; Pawlicka et al., 2004; Reiter et al., 2009b; Vondrák et al., 1999). Various designs of working devices, with the emphasis on all-solid-state multilayer configurations have been developed. The great interest of researchers during the last decade in the optical, electrical and mechanical properties of advanced solid polymers, as reflected by the large number of studies developed worldwide, is due to the fact that the most promising electrochromic devices use allsolid-state materials. Here, we will indicate briefly some electrolytes used in electrochromic devices. Most attempts to obtain a highly conductive polymer electrolyte have been devoted to polyether-based electrolytes: PEO, PPO and alkali salt (PEO–H3PO4) (Araújo et al., 1995; Armand, 1983; Daniel, 1986; Martins and Sequeira, 1991; Plancha et al., 1995; Watanabe and Ogata, 1986). The values obtained for the conductivity are in the order of 10−4 S cm−1. At room temperature it is important for the electrochromic device to have an electrolyte that is not only highly conductive but also highly stable. At present, few data are available in this direction. Another important class of polymeric conductors involves solid-state polymeric proton transport electrolytes such as Nafion and related materials (Randin, 1982), which possess free sulphonic acid groups and high molecular weight polyethyleneimines (PEIs) complexed with simple protonic acids (Akhtar et al., 1988a,b). Gel electrolyte lies between liquid and solid electrolytes. It consists of a solvent, a supporting electrolyte and a polymer. For example, 30 mol.% of LiClO4 in PVDF and 30 mol.% of LiClO4 in PMMA present conductivities of 10−6 and 3 × 10−11 S cm−1, respectively, at 25 °C (Tsuchida, 1983). Gel electrolytes have a good performance, such as high conductivity and adhesion, owing to enhanced viscosity. However, the polymer dissolved in a gel electrolyte is not a polymer matrix but a gelatising agent. A gel becomes liquid at elevated temperature and solidifies at lower temperature. The gel electrolytes introduced here could be regarded as an answer for large-scale electrochromic windows.
1.6
Conclusions
The past three decades have been devoted to establishing the majority of the polymer electrolytes and gathering understanding of their conduction mechanism, and their structure–conductivity relationships. In addition to solvating polymers, there are now clear indications that the materials should no longer be compared to solutions but to molten salts (Araújo et al., 1996b; Sequeira et al., 1998). The main driving force has been the improvement of the conductivity values, especially at room temperature and below, with the perspective of battery and fuel cell applications.
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The possibilities offered by polymer science are so wide that the scope of polymer electrolytes goes far beyond simple electrolytic media. For instance, a key feature of macromolecules is microphase separation on a nanometre scale and this can be considered as a simple case of a selfassembling structure. In the same vein, spatially defined microdomains with a given chemical or physical nature, hydrophilic/phobic, polar/non-polar, protic/aprotic, liquid crystalline, electronically conducting, etc., can be included in polymer electrolytes to form selective, multifunctional materials. The coexistence of ionic and electronic conductivity at the molecular level is being explored and will certainly prove to be fundamentally interesting and beneficial to exploit a wider choice of electrode reactions. Microdomains dispersed in the polymer matrix may serve as reservoirs for electrocatalysts, molecular recognition moieties, redox dyes, etc. Sensors, improved energy storage devices (batteries, supercapacitors), fuel cells, displays and photocells are just a few of the possible domains of application for compounds that tend more and more to mimic the properties of living matter.
1.7
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murphy tc and kramer we (1994), Proc Fourth Int Sem on Double Layer Capacitor and Similar Energy Storage Devices, Florida Educational Seminars, USA. nagasubramanian g, shen dh, surampudi s, wang q and surya prakash gk (1995), Electrochim Acta, 40, 2280. nakagawa h, iguchi s, kuwana k, nukuda t and aihara y (2003), J Electrochem Soc, 150, A695. neburchilov v, martin j, wang h and zhang j (2007), J Power Sources, 169, 221. neyertz s, brown d and thomas jo (1995), Electrochim Acta, 40, 2063 (1995). nguyen ca, xiong s, jan m, lu x and lee ps (2009), J Phys Chem B, 113, 8006. niedzicki l, kasprzyk m, kuziak k, zukoswka gz, armand m, bukowska m, marcinek m, szczeci ski p and wieczorek w (2009), J Power Sources, 192, 612. niitani t, shimada m, kawamura k and kanamura k (2005), J Power Sources, 146, 386. nishino a, yoshida a, tanahashi i, tajima t, yamashita m, muranaka t and yoneda h (1985), Nat Tech Rep, 31, 318. north jm, sequeira cac, hooper a and tofield bc (1982), Extended Abstracts, 1st International Meeting Lithium Batteries, April 1982, Rome. nozik aj (1978), Annu Rev Phys Chem, 29, 189. ouyang z, chen y, yao y and chen q (2007), Acta Polym Sin, 12, 1196. othman l, chew kw and osman z (2007), Ionics, 13, 337. paddison sj (2003), Annu Rev Mater Res, 33, 289. papke bl, ratner ma and shriver df (1981a), J Phys Chem Solids, 42, 493. papke bl, dupon r, ratner ma and shriver df (1981b), Solid State Ionics, 5, 685. papke bl, ratner ma and shriver df (1982), J Electrochem Soc, 129, 1694. passerini s, ressler jj, owens db and smyrl wh (1995), in Proc Symp Electrochemical Capacitors, Delnick FM and Tomkiewicz M (eds), Pennington, NJ, The Electrochemical Society, 86. patrick aj, glasse md, latham rj and linford rg (1986), Solid State Ionics, 18–19, 1063. patrick aj, glasse md, latham rj, sequeira cac and linford rg (1987), Abstracts 1st International Symposium Polymer Electrolytes, St Andrews, June, Paper 28. pawlicka a, dragunski dc, guimarães kv and avellaneda co (2004), Mol Cryst Liq Cryst, 416, 105. payne va, xu jh, forsyth m, ratner ma, shriver df and de leeuw sw (1995), Electrochim Acta, 40, 2087. plancha mjc, rangel cm and sequeira cac (1991), Extended Abstracts, 179th Meeting Electrochem Soc, Washington, DC, May 5–10, 1991, Vol 91-1, paper 801, 1142. plancha mjc, rangel cm and sequeira cac (1992a), Solid State Ionics, 58, 3. plancha mjc, rangel cm and sequeira cac (1992b), 8th Nat Conf of Physics 92, III P73, Vila Real, UTAD, 275. plancha mjc, rangel cm and sequeira cac (1992c), Extended Abstracts, 2nd Int Symp Electrochem Impedance Spectroscopy, Santa Barbara, CA, July 12–17. plancha mjc, rangel cm, sequeira cac and hudson mj (1993a), Extended Abstracts, Electrochem Soc Spring Meeting, Vol 93-1, 2455. plancha mjc, rangel cm, sequeira cac and hudson mj (1993b), Journées d’Electrochimie 93, Grenoble, 7–10 June, Abstract CA 1-9. plancha mjc, rangel cm and sequeira cac (1994), Extended Abstracts, 45th Annual Meeting Int Soc Electrochem, Porto, Vol 2, V-63.
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plancha mjc, rangel cm and sequeira cac (1995), Proc Seminar on Modern Applications of Polymeric Materials, University of Napier, UK, 118. plancha mjc, rangel cm and sequeira cac (1997a), Port Electrochim Acta, 15, 121. plancha mjc, rangel cm, sequeira cac and marquis fds (1997b), Técnica, 1/97, 89. plancha mjc, rangel cm, sequeira cac and hudson mj (1997c), J Appl Electrochem, 27, 1290. plancha mjc, rangel cm and sequeira cac (1998), J Electroanal Chem, 442, 91. plancha mjc, rangel cm and sequeira cac (1999), Solid State Ionics, 116, 293. ponce de leon c, walsh fc, rose a, lakeman jb, browning dj and reeve rw (2007), J Power Sources, 164, 441. poulsen fh, anderson nh, clausen k, skaarup s and sorensen ot (eds) (1985), Transport Structure Relations in Fast Ion and Mixed Conductors, Roskilde, Riso National Laboratory. prater k (1990), J Power Sources, 29, 239. raducha d, wieczorek w, such k, florjanczyk z and stevens jr (1996), J Phys Chem, 100, 20126. rajendran s, ramesh prabhu m and usha rani m (2008), J Appl Polym Sci, 110, 2802. randin jp (1982), J Electrochem Soc, 129, 1215. ratner ma (1987), in Polymer Electrolyte Reviews – 1, MacCallum and Vincent CA (eds), London, Elsevier, 173. reiche a, steurich t, sandner b, lobitz p and fleischer g (1995), Electrochim Acta, 40, 2153. reiter j, dominko r, nádherná m and jakubec i (2009a), J Power Sources, 189, 133. reiter j, krejza o and sedlariková m (2009b), Sol Energy Mater Sol Cells, 93, 249. riedinger h and faul w (1988), J Membrane Sci, 36, 5. robitaille cd and fauteux d (1986), J Electrochem Soc, 133, 315. saito y, stephen am and kataoka h (2003), Solid State Ionics, 160, 149. samms sr, wasmus s and savinell rf (1996), J Electrochem Soc, 143, 1225. sammels af and ang pgp (1984), J Electrochem Soc, 131, 617. sammels af and schmidt sk (1985), J Electrochem Soc, 135, 520. santos dmf (2006), MSc Thesis, Instituto Superior Técnico, TU Lisbon, Portugal. santos dmf (2009), PhD Thesis, Instituto Superior Técnico, TU Lisbon, Portugal. santos dmf and sequeira cac (2006), Defect Diffus Forum, 258–260, 333. santos dmf and sequeira cac (2007), ECS Trans, 3(32), 69. santos dmf and sequeira cac (2009a), J Electrochem Soc, 156, F67. santos dmf and sequeira cac (2009b), J Electroanal Chem, 627, 1. santos dmf and sequeira cac (2010), J Electrochem Soc, 157, B13. santos dmf, condeço jad, franco maw and sequeira cac (2007), ECS Trans, 3(32), 19. savadogo o and roberge pr (eds) (1997), Proc Second Int Symposium on New Materials for Fuel Cell and Modern Battery Systems, July 6–10, Montréal. savadogo o, roberge pr and veziroglu tn (eds) (1995), Proc First Int Symposium on New Materials for Fuel Cell Systems, July 9–13, Montréal. schantz s and torell lm (1993), Solid State Ionics, 60, 47. schmutz c, tarascon jm, gozdz as, warren pc and shokoohi fk (1994), Electrochem Soc Proc, 28, 330. seki s, kobayashi y, miyashiro h, mita y and iwahori t (2005), Chem Mater, 17, 2041. sequeira cac (1983a), Abstracts, 6th Annual Meeting Port Chem Soc, Aveiro, October 18–21, paper PB65.
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2 Ceramic polymer electrolytes J. S. S Y Z D E K, Warsaw University of Technology, Poland
Abstract: In this chapter a brief review on solid polymeric electrolytes is given, followed by the discussion of challenges and requirements to fulfil in order to commercialise them. Among various ways of polymeric electrolyte modification, composites with ceramic phases are discussed in detail here. Special attention is paid to electrolytes for lithium cells, based on polyethers. The main achievements, theories developed and precautions for future researchers willing to work in this area are discussed. Key words: composite polymer electrolyte, lithium battery, all solid state battery, surface treatment, ceramic filler, polymer-in-ceramic composite.
2.1
Introduction
Ever since electrochemical cells became objects of interest not only to researchers but also to consumers, their development sped up a lot and forced significant changes in the materials and geometries used. The size, weight, resistance to mechanical and thermal shocks, flexibility of the operating conditions, price, lifetime, and safety to the user and the environment of cells (more commonly, but not very correctly, called ‘batteries’), started to rule this business very quickly. Apart from optimisation of electrode materials, which is not discussed here, a lot of effort was devoted to the development of electrolyte systems, tailored to the specific electrochemical system. The role of electrolyte is two- or sometimes threefold: • •
•
It should provide ionic contact between electrodes, allowing the circuit to be closed when the cell is operational. It should assure electronic and spatial separation of the positive and negative electrode in order to avoid short-circuit and as a result selfdischarge of the cell, which in some cases can be very spectacular (as those of failed high power Li-ion cells). In the case of electrochemical systems where electrode components are not the only reactants appearing in the overall cell reaction, the electrolyte is the source (storage) of the remaining ones:
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(pos.) (neg.) (pos.) Pb(neg.) + PbO2(s) + 2H 2 SO4(el.) PbSO4(s) + PbSO4(s) + 2H 2 O(el.) (s)
In most cases the third function of the electrolyte is not required (or very limited), therefore from the point of view of the cell mass balance, the electrolyte is considered as a dead mass which should be cut down to the indispensable minimum. Taking into account simple Zn–MnO2 cells, the change that took place between the first and commercial design is more than clear when looking at the geometry evolution which allowed the amount of the electrolyte to be decreased dramatically (Fig. 2.1). In general, electronic and spatial separation can be obtained easily with application of dielectric solids. On the other hand solutions of proper electrolyte (acid, base or salt) in a suitable liquid solvent can provide good ionic conductivity needed for operating the cell. For many years these two functions have been separated and provided by different phases, i.e. liquid electrolytes soaked into a porous separator. In the 1970s M.B. Armand suggested that certain polymers (for example poly(oxyethylene)) that are characterised by a dielectric constant and donor number great enough to dissolve metal salts could be used as electrolytes in electrochemical cells provided that the ionic conductivity is high enough (Armand 1978; Armand et al. 1978, 1979). The history of this discovery has been nicely described in a review paper by P.V. Wright (Wright 1998) who was the first to report on the structure and ionic conductivity of poly(oxyethylene) complexes with alkali metal salts back in 1973 (Fenton et al. 1973; Wright 1998). However, the first reports on interactions of salts with poly(oxyethylene) are even older (Lundberg et al. 1966). The advantages of using solid polymer electrolytes in commercial lithium cells (up-to-date systems of choice for mobile applications) could be numerous (Tarascon and Armand 2001). They could offer: • • • •
•
• • • •
non-volatility; no decomposition at the electrodes; no possibility of leaks; use of metallic lithium in secondary cells (lithium dendrites growing on the electrode surface would be stopped by the non-porous and solid electrolyte); lowering the cell price (poly(ethylene oxide), PEO, is cheaper than organic carbonates; it could be used as a binder for electrodes to improve the compatibility of consecutive layers; moreover fabrication of such a cell would be easier – hence cheaper); strengthening of cells thanks to the all-solid-state construction; shape flexibility; lowering the cell weight – non-volatile, all-solid-state cells do not need heavy steel casing; improved shock resistance;
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2.1 Evolution of electrochemical cell design. Notice the decrease of the volume of electrolyte and the thickness of the separator leading to superior energy and power density. © Woodhead Publishing Limited, 2010
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better overheat and overcharge allowance; improved safety.
To take advantage of all these features, low interfacial resistances, high Li+ transport number and most of all high enough conductivity at ambient and subambient temperatures are required (Bruce and Vincent 1990; Gray 1991). These complexes were very soon recognised to be good ionic conductors but only above the melting point of the salt–polymer complex crystalline phase (Wright 1975). This is a result of the strong tendency of poly (oxyethylene) to crystallise, and crystalline phases are characterised by much lower conductivities than the amorphous ones (Lee and Wright 1982; Payne and Wright 1982; Wieczorek et al. 1988). A lot of effort was devoted to improve conductivity of these materials worldwide. The different approaches are as follows: •
Preparing electrolytes based on different polymer architectures (Fig. 2.2) (polyphosphazene, Blonsky et al. 1984; poly(siloxane), Lyons et al. 1996; poly(oxypropylene)) whose crystallinity is much lower than that of PEO (Florjan´czyk et al. 1992b; Silva et al. 1996). Another approach is preparation of polymer blends and/or copolymers (Florian´czyk et al. 1991). However, other issues are to be considered here, namely stability of the polymer against low and high potential materials, dielectric constant and donor number that have to be high enough to enable salt dissociation • Preparing poorly crystalline polymers in which sequences of oxyethylene units are longer than 4 and shorter than 15. Such structures are long enough to provide efficient solvation of lithium cations and short enough to avoid crystallisation. Such structures can be obtained by preparing copolymers with oxymethylene and/or oxypropylene units spread along the polymeric chain built of oxyethylene units. They can be also grafted or comb-like copolymers in which oligo(oxyethylene) units are attached to a main chain of high flexibility (for example poly(siloxane)), Florian´czyk et al. 1990, 1992a) • Doping the polymer matrix with light molecules, often of high dielectric constant and ability to dissolve the salt (organic carbonates, carboxylic esters, nitriles, glycols. . .). This plasticisation of the polymer leads to the increase in conductivity; however, it causes deterioration of the mechanical properties that are of main interest when using polymer electrolytes. In order to obtain more stable structures, such electrolytes are often reticulated by means of high energy radiation. In addition, plasticisers embedded in the electrolyte can lead to the problems encountered in ‘classical’ liquid electrolytes (evaporation, internal pressure, corrosion, flammability, etc.) (Girish Kumar and Munichandraiah 2000; Golodnitsky et al. 1996; Lee et al. 2000; Morita et al. 1997, 2000; Qian et al. 2002; Pitawala et al. 2008).
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Polymer electrolytes O
∗
∗
n Poly(oxyethylene)
O
∗
∗ n
Poly(oxypropylene)
∗
O
O 4–15
n
∗
Oxymethylene-linked-poly(oxyethylene) (OMPEO)
∗
O R1
∗
Si R2
n
Poly(siloxane)
∗
N R1
∗
P R2
n
Polyphosphazene
2.2 Polymer structures most commonly used as solid solvents for electrolytes.
•
Deeper studies on the polymeric electrolytes led to the conclusion that both charge carrier creation and migration can be limiting steps of the conduction phenomenon (Almond and West 1983; Siekierski and Wieczorek 1993), therefore not only the polymeric matrix determines the conductivity but also the dopant salt. Important aggregation of ions (forming ion pairs, ionic triplets and higher aggregates) (Dupon et al. 1982) decreases the number of charge carriers and leads to stiffening of the polymeric matrix due to bridging the coordination centres through the aggregates. This causes a drop in conductivity, combined with a drop in the cation transference number. In order to avoid that, new salts of very low lattice energy were developed and applied in electrolytes. As
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anions, very weak Lewis bases were applied (trifluoromethanesulphonate, bis(trifluoromethylsulphonyl) imide, etc. Such salts, often having a plasticising effect on the polymer matrix, form eutectics with low melting temperatures leading to highly conductive materials (Armand and El Kadiri 1987; Armand and El Kadiri Cherkaoui el Moursli 1985a,b,c). However the stability of such structures in a commercial cell is often an issue because of corrosion and stability problems (Behl and Plichta 1998; Chen et al. 1999; Krause et al. 1997). The last but not the least approach is addition of ceramic fillers. There is a variety of ceramic materials that offer relatively high ionic conductivities at room temperature (Goodenough et al. 1976; Hong 1976; Hooper 1977; Kafalas and Hong 1978; Sebastian and Gopalakrishnan 2003), however, applying them directly as electrolytes in commercial cells for mobile applications is not possible because they are brittle, difficult to process, and they cannot provide contact with the entire surface of porous electrodes made of powders as easily as polymers and liquids can do. Thus the first approach was using the polymeric electrolytes as conductive binders for the ceramic ionic conductors added to the polymer in a form of powder. Then inert materials were used. This will be described in detail in the next section.
2.2
Experimental approaches
2.2.1 Preparation of composites Dry polymeric electrolytes consist of a salt (or mixture of salts) dissolved directly into the polymer, without addition of any solvent or low molecular weight plasticiser. This feature distinguishes them from so-called gel– polymer electrolytes where the addition of a solvent is common and necessary to promote ionic transport. However, in order to prepare a polymeric electrolyte some solvents are usually applied. Owing to extremely high viscosities of high molecular weight polymers, even far above their melting point, dissolution of salts directly into the polymer can take months or years. Therefore in most cases such electrolytes are prepared via a casting technique. Polymer is first dissolved into a proper solvent, then the salt is added, the mixture is stirred overnight and then casted onto a glass, or preferably a Teflon® substrate (PEO is adhesive which makes the removal of electrolyte difficult if the substrate shows affinity to the polymer which is the case of glass). Once the solvent has been evaporated at atmospheric pressure and below solvent boiling point (evaporation that is too fast leads to the formation of bubbles/inhomogeneities within the sample) then vacuum drying is applied, often combined with elevated temperatures (Ravn Sørensen and Jacobsen 1982; Syzdek et al. 2007).
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(a)
(b)
2.3 Agglomeration and sedimentation of the filler during composite preparation. The upper and lower sides of the same sample exhibit large difference in filler loading (a) although the primary particle size is in the range of nanometers (40 nm). The next pair shows a sample prepared with adequate care. Both sides are identical despite the coarse character (2 µm) of the filler (b).
In general, composite electrolytes are obtained when filler is added to the polymer–salt solution in the solvent before the casting step. The homogeneity of such a composite slurry as well as that of the final solid composite is often questionable. The size of filler grains in the final material is also much larger than the initial one; these issues will be discussed later (Fig. 2.3). Another way of preparation of composites is the ‘hot pressing’ technique (Appetecchi et al. 2003). In short, electrolyte components (polymer, salt and
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the filler) are cooled to liquid nitrogen temperature, and then ball milled in order to mix them well and assure homogeneity of the mixture. It is subsequently pressed in a mould at a temperature above the polymer melting point and under relatively high pressure. High energy ball milling, however, can affect the molecular weight of some polymers, therefore applicability of this technique must be judged experimentally.
2.2.2 Examination methods A complete description of the examination methods of polymeric electrolytes is given in a separate chapter. Just a brief reminder is given here in order to help the reader follow this chapter. The most common method of electrolytes examination is electrochemical impedance spectroscopy (EIS) with blocking (usually stainless steel) and/ or reversible electrodes (most often metallic lithium, or other metal, common with the salt used; Saraswat et al. 1989). It allows determination of samples’ conductivity (Bauerle 1969), following the evolution of complete and half cells resistance upon storage (Sannier et al. 2007; Syzdek et al. 2007); sometimes it is used for transference number measurements (Ravn Sørensen and Jacobsen 1982), as well as a complementary technique in the studies on the crystallisation of solid electrolytes (Marzantowicz et al. 2006a,b,c, 2008), where detailed equivalent circuit considerations are applied. Another very important way of studying these systems is a set of DC techniques, i.e. voltammetry (Armand et al. 1980) and DC polarisation (chronoamperometry and chronopotentiometry). They allowed studies of the electrode reactions and examination of transport properties of electrolytes, especially transference number by Bruce and Vincent and co-workers (Bruce and Vincent 1987, 1990; Bruce et al. 1987; Christie et al. 1999; Evans et al. 1987; MacCallum et al. 1986) and/or Newman method (Doyle and Newman 1995; Hafezi and Newman 2000). Apart from electrochemical characterisation, a set of spectroscopic techniques is very useful. These are infrared (IR) (Marcinek et al. 2004), Raman (Marcinek et al. 2005) and nuclear magnetic resonance (NMR) (Berthier et al. 1983; Dai et al. 1998) spectroscopies. These allow observation of the interactions among electrolyte components, as well as providing chemical information about their structure. Pulse field gradient (PFG) NMR (Gorecki et al. 1988) is also used for observation of mass transport properties. Structural data can be obtained by X-ray diffraction (XRD) (Fenton et al. 1973; Lightfoot et al. 1992, 1994; MacGlashan et al. 1999; Pradhan et al. 2005). When combined with differential scanning calorimetry (DSC) (Gorecki et al. 1988; Sannier et al. 2007; Wieczorek et al. 1995) it can give insight into the crystallinity of the material. In addition positron annihilation lifetime spectroscopy (PALS) can give us information on the properties
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of the amorphous phase contained within the electrolyte sample (Asaad et al. 2008; Forsyth et al. 2002; Pas et al. 2005). Scanning electron microscopy, together with X-ray energy dispersive spectroscopy (EDS) and Raman microscopy can provide information on the morphology and homogeneity of the material (Syzdek et al. 2007; Wieczorek et al. 1995). Some caution must be kept in mind however. Polymers are not liquids, therefore most ‘classical’ techniques are not easily applicable, for example NMR spectroscopy, where broadening of peaks requires use of solid state (magic angle spinning) NMR, transference number determination by Hittorf–Tubandt (Saraswat et al. 1989) or the concentration cell method (Bouridah et al. 1985), which requires very special cell design and, because of slow diffusion in solids, takes a lot of time. On the other hand these are very soft materials that can be easily damaged upon examination with some ‘solid state’ techniques. Scanning and transmission electron microscopies (SEM, TEM) with a high energy electron beam can melt or even ‘burn’ the material (Syzdek et al. 2007). Large doses of X-ray radiation (X-ray diffraction XRD, X-ray absorbtion spectcoscopy XAS) or exposure to a laser beam (as in Raman spectroscopy) can damage the material as well. Furthermore the multiphase character of these materials makes it difficult to interpret all the experimental data.
2.3
First composites – conductive fillers
The technological limit of ionic conductivity of an electrolyte lies around 10−4 S/cm (Tarascon and Armand 2001). It can be lower, provided that one is able to prepare very thin electrolytes with sufficient mechanical strength. The first reported conductivities of dry PEO-based electrolytes were much lower at room temperatures, and reached the desired values only around and above the melting point of the PEO melting point (Wright 1975). On the other hand several ceramic ionic conductors were known (NASICON, β-alumina, inorganic glasses), but could not be applied in cells as electrolytes due to their brittle nature and difficulties with wetting the entire mass of the electrodes’ active material (Goodenough et al. 1976; Hong 1976; Hooper 1977; Kafalas and Hong 1978; Sebastian and Gopalakrishnan 2003). The first idea was to fill a polymeric electrolyte with powdered ceramic ionic conductors of high conductivity (Płocharski et al. 1989; Przyłuski and Wieczorek 1989; Wieczorek et al. 1989). The role of polymer–salt complex was to act as a conductive binder. It could also provide ionic conductivity within the active material thanks to very low distances between grains and large surface area. In the case of NASICON, very poor reproducibility of conductivity results was observed, and the reliable results were for low filler loadings (<5%v/v). For greater loadings impedance spectra became more
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complex than expected, exhibiting two relaxation phenomena. One relaxation phenomenon can be assigned to originate from the bulk conductivity, and the other from the interface between the polymeric matrix and the filler grains. The interfacial resistance turned out to be much larger than the bulk, and as such will limit considerably the enhancement of the conductivity by the conductive additive. Other attempts employing β- and β″-alumina as well as RbAg4I5 and KAg4I5 revealed high conductivities but only when the loading exceeded ca. 85%v/v, which is much more than the percolation threshold (Skaarup et al. 1988; Stevens and Mellander 1986). However, even these conductivities were much lower than those of the ceramic phase itself. On the other hand some enhancement of the conductivity at low filler loadings was observed and it could not be attributed to the percolation through the dispersed conductive filler phase. So what was the reason for the limited, but observed enhancement of the conductivity? NMR studies revealed that in the case of composite electrolytes the fraction of the crystalline phase is significantly lower than in the case of simple salt–polymer complexes (Płocharski et al. 1989; Wieczorek et al. 1989). In general, the composition of polymer–salt complexes is a result of rather complicated equilibrium and some non-equilibrium, kinetically limited phenomena (Fauteux and Robitaille 1985; Fauteux et al. 1985; Lee and Crist 1986; Minier et al. 1984; Munshi and Owens 1986; Robitaille and Fauteux 1986; Stainer et al. 1984). According to the Gibbs phase rule (Gibbs 1870; Mindel 1962), for all the compositions ranging from pure polymer up to that of the ‘thinnest’ crystalline complex, two phases should be present: pure, crystalline PEO and pure PEO–salt crystalline complex. Nevertheless, polymeric materials are ‘intrinsically impure’, for example due to their polydispersity. Additionally, their crystallisation is kinetically limited, therefore in all polymeric materials there are always amorphous domains. Thus PEO–salt complexes usually consist of three phases (Fig. 2.4) – pure crystalline PEO, crystalline PEO–salt complex and amorphous PEO-salt complex; the latter is of undefined composition (Wieczorek et al. 1989). The very first XRD-based results by P.V. Wright concluded that the conductivity of salt–polymer complexes decreases with an increase of their crystallinity (Fenton et al. 1973). This was confirmed later by Berthier et al. (1983). Coming back to the conductive filler problem – since their addition decreases the crystallinity (as confirmed by NMR), it is not surprising that the conductivity of the electrolytes is enhanced upon their addition. On the other hand, the presence of polymer–filler interfaces characterised by very high impedance prevents one from taking advantage of ionic conductivity of the filler itself. Therefore all we gain upon filler addition is the amorphitisation of the polymer–salt complex itself (Płocharski et al. 1989; Wieczorek et al. 1990).
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Polymer electrolytes 100 40
20
12
8
6
4
3
150
Temperature (˚C)
L
L + P(EO3 · LiClO4)
100
P(EO6 · LiClO4) + P(EO3 · LiClO4)
50
PEO + P(EO6 · LiClO4)
0
0.1
0.2
0.3
0.4
XLiClO4 (wt)
2.4 A phase diagram of the PEO-LiClO4 system. (䊉) – transitions observed by microscopy, (+) – transitions observed by conductivity measurements. Reproduced with permission from Robitaille and Fauteux J. Electrochem. Soc. 133, 320. Copyright 1986, The Electrochemical Society.
2.4
Development of insulating fillers
Researchers soon found that it is very difficult to take advantage of the fillers’ conductivity. It is possible only at very high loadings and at that point the mechanical integrity of the material is highly questionable. The advantage of filler addition at low loadings is the amorphitisation. But why should the polymer be less crystalline with filler added? It is not difficult to understand when we think what happens upon cooling the polymer melt. At the melting point crystallisation begins at each point where a crystallisation
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nucleus is present. Filler grains can be considered as nucleation centres. Their concentration is much larger than those of nuclei intrinsically present in a polymer matrix. Therefore crystallisation starts simultaneously in very many points. Growing spherulites meet quickly, leaving a large amount of amorphous phase in between. Since the first recognised function of fillers was their ability to promote fast and incomplete crystallisation of the polymeric matrix, researchers decided to try simpler and cheaper fillers, for example alumina. It does not conduct either electrons or ions, but this will not worry us because high interfacial impedances would not allow us to take advantage of filler conductivity anyway (Wieczorek et al. 1989). The very first attempt to add alumina to PEO–LiClO4 complex was driven by the hope of improving the mechanical strength of this material (Weston and Steele 1982) (filling construction polymers with inorganic powders in order to gain mechanical strength and decrease their price was already common then; De Gennes 1971). That was accomplished with ca. 40 µm α-Al2O3 powder. Further studies revealed that by using finer grains of γ-Al2O3 (0.3–1 µm) an increase in ambient temperature ionic conductivity was achieved together with strengthening the material (He et al. 1986). The differences between the samples filled with alumina powder of the same structure but just different grain size inspired other research groups to examine in detail the influence of alumina powders on the conductivity of composite polymeric electrolytes. Both loading and filler grain size were examined in detail in (PEO)10NaI–Al2O3 (α, ε and θ) system (the notation (PEO)xCA means that for each salt molecular formula CA (cation–anion) there are x oxyethylene units in the polymeric chain) (Płocharski et al. 1989; Przyłuski et al. 1990; Wieczorek et al. 1989, 1990). The conclusion was drawn that by decreasing the filler particle size, an increase of the conductivity is obtained. The reason for the enhancement of transport properties was the decrease of the crystallinity index as confirmed by NMR, DSC, Raman spectroscopy as well as XRD. It should be kept in mind that on decreasing the particle size, for constant loading, the amount of grains in the unit volume increases. Therefore it was not surprising that on increasing the loading, the crystallinity index was also confirmed to decrease, regardless of the particle size. This supported the hypothesis of the fast and incomplete crystallisation of the electrolyte induced by the filler grains which leads to more amorphous materials, which in turn offer better transport properties (diffusion of ions as well as neutral permeates is much faster in amorphous polymers than in the crystalline ones). A very important feature is the fact that this amorphitisation is not just a temporary effect. Long-term measurements revealed no recrystallisation of electrolytes upon prolonged storage (Fig. 2.5). Although the crystallinity index decreases with increase of the filler loading, the conductivity profiles exhibit a maximum at certain filler
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Crystallinity index (%)
75 65 55 45 35 25 0
15 20 10 Filler volume fraction (% )
5
25
2.5 A decrease of the crystallinity of polymer electrolytes on addition of ceramic fillers (1M LiClO4 in 5 × 106 g/mol poly(oxyethylene) mixed with 5 µm TiO2).
Conductivity (S/cm)
1E-05
1E-06
1E-07
1E-08 0
5
10
15
20
25
Filler volume fraction (%)
2.6 Conductivity of 1M LiClO4 in 5 × 106 g/mol poly(oxyethylene) mixed with 5 µm TiO2 as a function of filler volume fraction. Data collected at 300 K.
concentration. This is a result of two factors. Firstly, increasing the amount of insulating phase, the dilution effect occurs. Secondly, closely arranged grains stiffen the polymeric matrix, restricting the segmental motion of the chains. Since movement of ions is in most cases coupled with it – its restriction means limiting the ionic conductivity. The loading for which optimum is observed lies around 3–6% v/v, which means 10–20% w/w for most of the ceramic materials used (Dissanayake et al. 2003). However, the position of this maximum depends strongly on the filler particle size. This is discussed in detail in papers devoted to modelling of composite systems, since there are very many different models describing the conductivity of composite systems. An example of such a study is given in Wieczorek et al. (1996d) (Fig. 2.6). © Woodhead Publishing Limited, 2010
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75
Impact of the filler surface on the transport properties
The fact that the conductivity of composites based on crystalline polymers is enhanced below their melting point appeared to be rather easy to explain, although not without a discussion on the transport mechanism (amorphous vs. crystalline phase). But then a new phenomenon was discovered. Conductivities of electrolytes based on very poorly crystalline polymers (OMPEO – oxymethylene-linked-poly(oxyethylene)) were observed to increase on filler addition although their very low crystallinity could not be lowered (Best et al. 1999; Johansson et al. 2001). Furthermore even in the case of easily crystallising polymers, conductivities measured above the crystalline phase melting point appeared to be higher in the case of composites, especially when the salt concentration was high (above that corresponding to the maximum conductivity for non-composite systems) (Capuano et al. 1991; Croce et al. 2006; Golodnitsky et al. 1997; Villarreal et al. 1999). What was also observed, was the increase of the cation transference number (Appetecchi et al. 2003; Croce et al. 2006). This observation led to the idea that the presence of filler alters interaction between ions and the polymeric matrix (Stevens and Wieczorek 1996; van Eijck et al. 2005; Wieczorek et al. 1996a,b; Zalewska et al. 1996). A well-known formalism developed by Fuoss and Kraus (Fuoss and Kraus 1933a,b,c; Kraus and Fuoss 1933) could shed some light on this, provided that viscosities of electrolytes could be measured. Viscosities of high molecular weight polymers, especially doped with salts are extremely high and as such difficult to measure; therefore an idea of using ethylene oxide oligomers as model systems emerged and took its place in the field of polymeric electrolytes (Wieczorek et al. 1998a). Experiments performed revealed that addition of salt increases viscosity of electrolytes (Marcinek et al. 2004, 2005). That is clear when one thinks about the salt as consisting of Lewis acid (cation, usually Li+ which is a hard and strong acid) and a Lewis base (anion, usually weak and soft bases are used). On the other hand the poly(oxyethylene) chains contain oxygen atoms, being rather strong and hard Lewis bases. Adding a small amount of salt into the polymer results in its dissociation and creation of cations solvated by the polymer oxygen atoms and anions. In order to move across the electrolyte, cations have to exchange some oxygen atoms in the coordination sphere into the new ones, coming from another part of the same chain, or one coming from another chain (intra- and inter-chain hopping) (Berthier et al. 1983). The kinetics of this is determined by the ligand exchange rates (Helm and Merbach 1999). Anions are more free to move in the polymer under electric field because they are not coordinated to the polymer chains, therefore cation transference numbers are usually much lower than 0.5 (Bruce and Vincent 1989; Doyle and Newman 1995; Evans et al. 1987; Fauteux et al. 1985).
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When the salt concentration increases, the equilibrium between free ions and ion pairs produces more and more of the latter, and when even higher concentrations are achieved, ionic triplets and higher aggregates appear. Such products together with ‘free’ lithium ions act as temporary bridges between oxygen atoms in the polymeric matrix producing the so-called ‘stiffening’ effect which lowers the conductivity (because movement of ions is in most cases coupled with the segmental motions of polymeric chains) (de Leeuw et al. 2001; van Zon and de Leeuw 2001). This effect is difficult to observe in solid systems, but is easy to follow in case of liquid at room temperature, oligo(oxyethylene)-based electrolytes by viscosity measurements. Afterwards, studies with composite systems were performed. A decrease of the viscosity was observed upon addition of Al2O3 combined with an increase of the cation transference number (Marcinek et al. 2005). This effect could result first of all from the interactions between the filler surface and other electrolyte components. Thus several different fillers were examined, with carefully designed surface groups (Ciosek et al. 2005). Alumina with neutral, basic and acidic (both Lewis type) surfaces was introduced into PEOME–LiClO4 electrolytes of various concentrations and examined with a variety of techniques. At higher salt concentrations when aggregation of ions occurs, fillers led not only to a decrease of the viscosity, but also to an increase in the fraction of free ions. The best results were obtained in the case of acidic fillers (Fig. 2.7). This results from interaction of acidic surface states with anions (Lewis bases) present in electrolyte as free ions, ion pairs, triplets, etc. The latter especially affect the viscosity and cation transference number. With the use of acidic filler, the fraction of free ions was increased as confirmed by Fuoss–Kraus calculations, as well as by Fourier transform infrared (FT-IR) investigations (Marcinek et al. 2005). The increase in the fraction of free ions means an increase of the number O
O S
O
O O
O
O O
O
Brønsted acid centre
S
Lewis acid centre
O
O
O
O
O
O
SO3H
O O
O
O
O
O
OH O
O
O
O
O
O
Me
Me
Me
Me
Me
Me
Me
Me
Me
O
O
O
O
O Crystal
O
O
O
O
2.7 Introduction of sulphate groups onto the surface of metal oxides as a way to increase their acidity (electron withdrawing effect of attached sulphate groups increases the acidity of metal cation Lewis acid centres; their monodentate attachment to the ceramic grain leads to formation of Brønsted acid sites).
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of charge carriers. Combined with lower viscosity (which means higher mobility of polymeric chains) both factors lead to an increase in conductivity. Similar investigations were conducted on high molecular weight PEO filled with both ceramic and organic fillers. In general, addition of AlCl3 (soluble) or Al2O3 (insoluble) with an acidic surface appeared to enhance both conductivity and cation transference number of the electrolyte (Appetecchi et al. 2003; Borkowska et al. 2001; Ciosek et al. 2005; Florianczyk and Wieczorek 1996; Johansson and Jacobsson 2004; Marcinek et al. 2000a, 2000b, 2004, 2005; Moskwiak et al. 2006; Quartarone et al. 1998; Siekierski et al. 1998; Stevens and Wieczorek 1996; Stygar et al. 2005; S´wierczyn´ski et al. 2001; Wieczorek 1995; Wieczorek et al. 1996b,c, 1998a,b). Some studies revealed that the enhancement of conductivity does not have to result from short-range interactions and changes of polymer conformation. The long range coulombic forces (and their alteration by the presence of filler) can be responsible for the changes in conductivity, viscosity and spectra of the electrolytes (Kearley et al. 2004). Since acidic fillers appeared to be effective in enhancing the transport properties, this topic was studied in further detail. In order to maximise the interactions, very strongly acidic surface groups were introduced onto the powder surface (Sannier et al. 2007; Syzdek et al. 2007). Such systems are still under investigation in order to determine if they are solid superacids (Marczewski et al. 2004). The Hammet numbers of these powders correspond very closely to that of 100% sulphuric acid. However, depending on the support either Brønsted or Lewis type surface groups can dominate. Using such very strongly acidic fillers very high values of cation transference number were obtained, increasing with the increase of the acid content, regardless the molecular weight of the polymer (both liquid and solid) (Syzdek et al. 2007) as listed in Table 2.1.
2.6
Interfacial concerns
The great success of acidic fillers provoked more detailed studies on these materials. The performance of a complete cell depends not only on the transport properties (conductivity and transference number) but also on the interfacial processes, i.e. on the kinetics of the charge transfer reactions (Egashira et al. 2006; Sannier et al. 2007) and on the nature of the products of the side reactions taking place at the interface, which sometimes lead to very poorly conducting phases. In classical liquid electrolytes the solid electrolyte interface (SEI) formation at potentials below 0.8 V vs. Li+/Li is commonly known (Morales and Acosta 1998; Munichandraiah et al. 1998; Peled 1983; Peled et al. 1995). These interfacial processes keep on going over time, causing deterioration of cell characteristics, both internal resistance, as well as capacity loss. In the case of solid electrolytes based on
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Table 2.1 Values of Li transference number obtained by the Bruce–Vincent method for polymer electrolytes consisting of 1M LiClO4 in PEO (5 × 106 g/mol) and 10% w/w of corresponding filler. Transport numbers measured by AC technique with Sørensen– Jackobsen formula lay within the range 0.28–0.41, with the same order. For noncomposite electrolyte it is 0.16. Data from the author’s own work Acid concentration w/w %
Type of support Coarse Al2O3
Micro Al2O3
Nano Al2O3
Coarse TiO2
Nano TiO2
Raw 0 1 4 8
0.47 0.53 0.68 0.71 0.72
0.42 0.69 0.72 0.75 0.79
0.72 0.57 0.63 0.69 0.73
0.63 0.61 0.66 0.72 0.77
0.72 0.73 0.77 0.81 0.86
Non-composite electrolyte
0.32
polyethers, which are much less corrosive towards lithium, the resistance of the interface with lithium metal also grows over time. Some models were proposed, explaining the observed increased impedances at the interface even when no reaction is expected to take place (Kerr et al. 2004). The positive effect of ceramic fillers on the nature and properties of the lithium–electrolyte interface was recognised quite early and the first suppositions were that the filler grains can gather at the interface, decreasing the surface area between the lithium electrode and the polyether. However that supposition was discarded with the help of experiments conducted with the use of both low and high molecular weight PEO-based electrolytes (Syzdek et al. 2007). Even more pronounced effects were observed when the ceramic filler was grafted with acid surface groups. In such cases the interfacial resistance was observed to grow slightly in the very first days and then to decrease to values lower than the initial one (Sannier et al. 2007; Syzdek et al. 2010). The resistance was observed to decrease slightly as the concentration of acid sites is increased. These long-term studies conducted at elevated temperatures confirmed that stability of such composites against lithium metal is true (Syzdek et al. 2007); however, the concentration of dopant salt can alter the situation significantly (Egashira et al. 2006; Bac et al. 2004, 2006a,b). Other studies showed that by adjusting the surface groups the efficiency and kinetics of lithium metal plating/stripping can be enhanced. Acidic fillers were observed to be the best choice (Egashira et al. 2006). How could that be explained? One supposition is that fillers of high Lewis acidity can trap impurities/traces of water/other reactive molecules (including salt anions) (Li et al. 2001) and in this way counteract secondary
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reactions at the lithium electrode. Another comes from observation, that greater stabilities were observed in the same samples for which the cation transference number was high (Bac et al. 2006a,b; Ciosek et al. 2005; Marcinek et al. 2005). Other authors emphasise that on the polarisation of the cell, whenever the transference number of the electroactive species is lower than unity, at one electrode there is depletion and at the other enrichment of the salt. Since the composition of the electrolyte is optimised from the viewpoint of bulk conductivity, any departure from the average concentration leads to formation of poorly conductive zones (Kerr et al. 2004; Munichandraiah et al. 1996). The last model described in the literature comes from the explanation according to the space charge effect and the fact that next to the interfaces, due to surface tensions, equilibrium compositions may differ from the bulk (average) ones, thus leading to zones of different conductivities (Mattsson et al. 2000). Whatever the model we assume, the experimental fact is that highly acidic fillers enhance stability and transport properties of the interfaces with low potential electrodes as presented in Fig. 2.8.
R/Ω 250
R el. + RCT + R passiv. + R diff R el. + RCT + R passiv. R el. + RCT R el.
200
150
100
50
0 0
5
10
15
20 t/day
25
30
35
2.8 Evolution of the resistance of a symmetrical lithium cell with a composite electrolyte consisting of 1M LiClO4 in 5 × 106 g/mol poly(oxyethylene) with 10% w/w 1 µm Al2O3 modified with sulphuric acid as a filler. Storage and measurements at 70 °C. R el. – resistance of the electrolyte, RCT – resistance of the charge transfer, R passiv. – resistance of the passive layer, R diff – diffusion resistance. Notice initial decrease followed by a stability period of the measured impedances. This behaviour is very different from the continuous increase observed for liquid electrolytes.
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2.7
Other types of ceramic–polymer systems
In most of the systems described so far, the bulk properties of the ceramic filler were of secondary importance. That was because of high impedances of the filler–polymer interface. It was shown that by exposing composites to vapours of plasticising agents the adhesion of the matrix to the filler may be improved and the conductivity increased (Nairn et al. 1999), but the plasticisation of the polymer can lead to deterioration of its mechanical properties, as well as provoke some side reactions at the interface with low potential electrodes. Therefore taking advantage of the transport properties of the filler is problematic. However, there are some other bulk properties that we can apply in order to ameliorate the composite. One of examples is the application of ferroelectric filler, i.e. BaTiO3. This material is of very high dielectric constant that alters the long-range interactions in the composite system. One has to remember that on decreasing the particle size one can go beyond the size of ferroelectric domains and in this way lose the main advantage of this material. When very finely divided it acts in a similar manner as other ceramic powders. When the size is large enough it can disturb both the crystallisation of the polymer and the dissociation of salt, leading to both increased conductivities and cation transport numbers (Capiglia et al. 2002; Li et al. 2001; Morita et al. 2006; Park et al. 2003; Sun et al. 2000; Yang et al. 2006). Also the introduction of electronically conducting fillers at loadings below the percolation threshold was done leading to enhanced transport properties (Bloise et al. 2000; Cheng et al. 2004; Franco et al. 1999; Kerr et al. 2003; Lavall et al. 2005). Calculations revealed that the enhancement takes its origin in a decrease of the activation energy of charge carriers’ migration (Muszynska et al. 2002). An attempt to introduce ‘macroscopic anions’ was made. When cation exchanging clays were introduced into the polymer, some lithium cations were liberated into the polymeric phase and contributed to ionic transport. Negatively charged clay particles, being of macroscopic size, could not move and in this way the lithium transport number could be significantly increased (Aranda and Ruiz-Hitzky 1992; Chen and Chang 2001; Kim and Park 2007; Kim et al. 2008; Krawiec et al. 1995; Riley et al. 2002, 2003; Sandí et al. 2005; Vaia et al. 1995; Walls et al. 2003). Some other attempts were made to modify filler surface states, namely both hydrophilic and hydrophobic moieties were introduced onto the fumed silica surface. Afterwards these silicas were applied as fillers for polymeric electrolytes based on PEODME250 and such systems were examined from the viewpoint of rheology and conductivity (Fan et al. 1998). It appeared that although rheology depends strongly on the nature of surface groups (what is not surprising in case of a filler with so large a
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specific surface), the conductivity seems not to depend on the type of surface group at all. When surface groups were compatible with the polymer matrix the filler grains could interact with both the matrix and other grains. Then it is not surprising that surface groups that resemble the polymeric matrix do not produce any preferential interactions with any of the electrolyte components. This is also the case of inert surface groups. When the surface groups are incompatible with the polymer, they interact with those coming from other grains. In such a way a 3D matrix made of fillers is created, affecting the rheology but not the conductivity since the structure of the ion conducting phase is not affected in the micro-scale. In the case of high molecular weight poly(oxyethylene) addition of silica with polymercompatible surface groups causes an increase of crystallinity, but at the same time the glass transition temperature is decreased, leading to an amorphous phase with enhanced transport properties. Improved stability and mechanical properties follow this increase of total conductivity (Fan et al. 2003). Some further explanation comes with the help of molecular dynamics (MD) simulations (Johansson and Jacobsson 2004; Kasemägi et al. 2004) (modelling is discussed in Chapter 8). It was found that the mobility of charge carriers in the close proximity of the filler grains is very limited; however, it increases quickly with increasing distance. Therefore one should think about filler grains and their closest neighbourhood more as traps rather than places of fast ionic transport, which seems to be enhanced at some distance away from the particle (Kasemägi et al. 2003, Wieczorek, 1995). Some extended works have been devoted to studies on the influence of highly acidic fillers (both soluble such as AlBr3 or AlCl3 and ceramic ones) on the properties of gel polymer electrolytes based on PVdF. It was found that the conductivity and cation transference numbers increase on the addition of these strong Lewis acids. Also the stability of the interface with lithium metal is enhanced; however, in the case of aluminium halides the reactions with organic carbonates can take place and affect the properties of such systems on prolonged storage (Stolarska et al. 2007; Walkowiak et al. 2006, 2007). Recent work in the field of ceramic polymeric electrolytes is the so-called ‘polymer-in-ceramic’ approach presented by Syzdek et al. (2009). It consists of introducing polymer–salt complexes into porous ceramic structures. In this manner we split the spatial and electronic separation and ionic conductivity into two phases again, but the difference is that ceramics are much more thermally stable, resistant to reactions with lithium and to puncturing with lithium dendrites that can appear in the cell. Also the wettability of ceramic materials is much better than that of polymers, therefore this approach can be considered as an alternative to the Bellcore® process that
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is used for the production of plastic lithium ion cells (Doyle et al. 1996; Tarascon and Armand 2001; Tarascon et al. 1996).
2.8
Conclusions
The impact of fillers on the ionic conductivity of solid PEO-based systems has been a subject of discussion for quite many years. The literature shows that for some systems the enhancement of conductivity is tremendous, although for some others it is very limited or there is even no enhancement of transport properties at all. Some systems are stable, some are not, although they differ only slightly. A deeper look into the subject, combined with extended laboratory practice prove that ceramic fillers work, but the extent to which it happens depends a lot on the preparation conditions. Polymer–ceramic composites are heterogeneous systems: that is their nature and therefore the order in which we mix the components matters. The same applies to the thermal history of the materials and of the final product. As anyone who has tried to prepare angel food cake knows, the order of mixing and thermal history can change. Also the preparation conditions can be of tremendous importance. Just by changing them, working with the same material all the time, we can obtain composites that differ a lot, concerning their structure, transport and mechanical properties, etc. (Syzdek et al. 2010; Volel et al. 2004). Although Lord Rutherford said: ‘If your experiment needs statistics, you ought to have done a better one’, usually very many results need to be averaged over an important population of samples in order to enable the researcher to draw a reasonable conclusion. This is not possible either with single measurements, or with a limited set of experimental techniques. Since even the configuration of the experimental setup can affect the results, great care must be taken at each step of the work in order not to produce meaningless numbers. The last remark is about the future prospects. Most of the text presented here deals with lithium cells and lithium electrolytes. However one has to keep in mind that most of the knowledge earned on these systems can be transferred easily (if not directly sometimes) into different ones. A good example could be sodium cells. If one day humankind faces a scarcity of lithium (http://www.meridian-int-res.com/Projects/Lithium_Problem_2. pdf) it will have to move towards sodium cells. However, the chemistry of lithium and sodium electrodes with liquid electrolytes is quite different. On the other hand all the examination methods and experimental setups developed for dry polymer electrolytes apply to lithium as well as sodium cells. There are also some common conclusions (West et al. 1989). At the time this book is being written, revival of research on sodium electrolytes and
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sodium cells is taking place and the lithium know-how regarding the polymer electrolyte technology is being successfully transferred into the sodium area. For low scale, high power applications lithium cells seem to be indispensable. But even though they work fine with liquid electrolytes, soon lithium– polymer technology will be explored more and more, especially in the context of electric vehicle batteries. Li-ion technology employing liquid electrolytes is susceptible to runaway overheating (depending on the system, the threshold lies between 70 and 120 °C) or overcharging, although they work perfectly at usual ambient temperatures. However, there are polymer batteries, which need to be kept at temperatures around 60 °C in order to operate properly; these are much safer in use. If one thinks about a real system, introduction of a thermal management system is mandatory for both systems (liquid and polymer). Liquid electrolyte cells need to be protected from both high and low temperatures (that can be reached at high charge– discharge rates as well as in hot areas in summer, or can be reached in winter). Polymer cells need only to be heated to around 60 °C; overheating is not really an issue in this case (they are capable of surviving at 200 °C). Simplicity and low energy consumption of such systems (insulation+heating) have been demonstrated, and promise to cause more rapid development of polymer cells for electric vehicle applications. As for the composite systems, research on both regular, and reversed phase composites is being developed. Addition of ceramic powders keeps the principles of the classical polymer technology and pays in improved mechanical stability. The polymer-in-ceramic systems offer extraordinary mechanical properties and great conductivities; however, dealing with materials that behave in many ways like ceramics is a challenge for the battery production technology. One way or another, this new approach is expected to gain in popularity and develop rapidly in the coming years.
2.9
Acknowledgements
The author acknowledges warmly Professors Władysław Wieczorek from Warsaw University of Technology and Michel Armand from Université de Picardie Jules Verne (cited many times in this chapter) for their help with the preparation of this manuscript. He also thanks the European Union for its support in the framework of European Social Fund through the Warsaw University of Technology Development Programme. Of course not all issues related to composite electrolytes could be covered in this chapter but the author hopes it will be a useful guide to this area of research. Readers are encouraged to study original papers, keeping in mind all factors ruling properties of such systems.
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2.10
References
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payne, d.r. and wright, p.v., 1982. Morphology and ionic conductivity of some lithium ion complexes with poly(ethylene oxide). Polymer, 23(5), 690–693. peled, e., 1983. Film forming reaction at the lithium/electrolyte interface. Journal of Power Sources, 9(3), 253–266. peled, e., golodnitsky, d., ardel, g. and eshkenazy, v., 1995. SEI model – application to lithium–polymer electrolyte batteries. Electrochimica Acta, 40(13–14), 2197–2204. pitawala, h.m.j.c., dissanayake, m.a.k.l., seneviratne, v.a., mellander, b. and albinson, i., 2008. Effect of plasticizers (EC or PC) on the ionic conductivity and thermal properties of the (PEO)9LiTf:Al2O3 nanocomposite polymer electrolyte system. Journal of Solid State Electrochemistry, 12(7–8), 783–789. płocharski, j., wieczorek, w., przyłuski, j. and such, k., 1989. Mixed solid electrolytes based on poly(ethylene oxide). Applied Physics A Solids and Surfaces, 49(1), 55–60. pradhan, d.k., samantaray, b.k., choudhary, r.n.p. and thakur, a.k., 2005. Effect of plasticizer on structure – property relationship in composite polymer electrolytes. Journal of Power Sources, 139(1–2), 384–393. przyłuski, j. and wieczorek, w., 1989. Increasing the conductivity of polymer solid electrolytes: a review. Solid State Ionics, 36(3–4), 165–169. przyłuski, j., such, k., wycis´lik, h., wieczorek, w. and floria czyk, z., 1990. PEObased polymer blends as materials for solid electrolytes. Synthetic Metals, 35(1–2), 241–247. qian, x., gu, n., cheng, z., yang, x., wang, e. and dong, s., 2002. Plasticizer effect on the ionic conductivity of PEO-based polymer electrolyte. Materials Chemistry and Physics, 74(1), 98–103. quartarone, e., mustarelli, p. and magistris, a., 1998. PEO-based composite polymer electrolytes. Solid State Ionics, 110(1–2), 1–14. ravn sørensen, p. and jacobsen, t., 1982. Conductivity, charge transfer and transport number – an ac-investigation of the polymer electrolyte LiSCN-poly(ethyleneoxide). Electrochimica Acta, 27(12), 1671–1675. riley, m., fedkiw, p.s. and khan, s.a., 2002. Transport properties of lithium hectoritebased composite electrolytes. Journal of the Electrochemical Society, 149(6), A667–A674. riley, m.w., fedkiw, p.s. and khan, s.a., 2003. Lithium hectorite clay as the ionic conductor in LiCoO2 cathodes. Journal of the Electrochemical Society, 150(7), A933–A941. robitaille, c.d. and fauteux, d., 1986. Phase diagrams and conductivity characterization of some PEO–LiX electrolytes. Journal of the Electrochemical Society, 133(2), 315–325. sandí, g., kizilel, r., carrado, k.a., fernández-saavedra, r. and castagnola, n., 2005. Effect of the silica precursor on the conductivity of hectorite-derived polymer nanocomposites. Electrochimica Acta, 50(19), 3891–3896. sannier, l., zalewska, a., wieczorek, w., marczewski, m. and marczewska, h., 2007. Impact of ‘Super Acid’ like filler on the properties of a PEGDME/LiClO4 system. Electrochimica Acta, 52(18), 5685–5689. saraswat, a.k., magistris, a., chiodelli, g. and ferloni, p., 1989. Electrochemical behaviour of the cell Pb/PEO40. Pb(ClO4)2/Pb. Electrochimica Acta, 34(12), 1745–1747.
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sebastian, l. and gopalakrishnan, j., 2003. Lithium ion mobility in metal oxides: A materials chemistry perspective. Journal of Materials Chemistry, 13(3), 433– 441. siekierski, m. and wieczorek, w., 1993. Application of the ‘universal power law’ to the studies of ac conductivity of polymeric electrolytes. Solid State Ionics, 60(1–3), 67–71. siekierski, m., wieczorek, w. and przyłuski, j., 1998. AC conductivity studies of composite polymeric electrolytes. Electrochimica Acta, 43(10–11), 1339–1342. silva, g.g., lemes, n.h.t., da fonseca, c.n.p. and de paoli, m., 1996. Solid state polymeric electrolytes based on poly(epichlorohydrin). Solid State Ionics, 93(1–2), 105–116. skaarup, s., west, k. and zachau-christiansen, b., 1988. Mixed phase solid electrolytes. Solid State Ionics, 28–30(Part 2), 975–978. stainer, m., hardy, l.c., whitmore, d.h. and shriver, d.f., 1984. Stoichiometry of formation and conductivity response of amorphous and crystalline complexes formed between poly(ethylene oxide) and ammonium salts: PEOx-NH4SCN and PEOx–NH4SO3CF3. Journal of the Electrochemical Society, 131(4), 784–790. stevens, j.r. and mellander, b., 1986. Poly(ethylene oxide)–alkali metal–silver halide salt systems with high ionic conductivity at room temperature. Solid State Ionics, 21(3), 203–206. stevens, j.r. and wieczorek, w., 1996. Ionically conducting polyether composites. Canadian Journal of Chemistry, 74(11), 2106–2113. stolarska, m., niedzicki, l., borkowska, r., zalewska, a. and wieczorek, w., 2007. Structure, transport properties and interfacial stability of PVdF/HFP electrolytes containing modified inorganic filler. Electrochimica Acta, 53(4), 1512–1517. stygar, j., zukowska, g. and wieczorek, w., 2005. Study of association in alkali metal perchlorate-poly(ethylene glycol) monomethyl ether solutions by FT-IR spectroscopy and conductivity measurements. Solid State Ionics, 176(35–36), 2645–2652. sun, h.y., takeda, y., imanishi, n., yamamoto, o. and sohn, h., 2000. Ferroelectric materials as a ceramic filler in solid composite polyethylene oxide-based electrolytes. Journal of the Electrochemical Society, 147(7), 2462–2467. s´wierczyn´ski, d., zalewska, a. and wieczorek, w., 2001. Composite polymeric electrolytes from the PEODME–LiClO4-SiO2 system. Chemistry of Materials, 13(5), 1560–1564. syzdek, j., borkowska, r., perzyna, k., tarascon, j.m. and wieczorek, w., 2007. Novel composite polymeric electrolytes with surface-modified inorganic fillers. Journal of Power Sources, 173(2), 712–720. syzdek, j., armand, m.b., marcinek, m., zalewska, a., z˙ukowska, g.z. and wieczorek, w., 2010. Detailed studies on the fillers modification and their influence on composite, PEO-based polymeric electrolytes. Electrochimica Acta, 55(4), 1314– 1322. syzdek, j., armand, m., gizowska, m., marcinek, m., sasim, e., szafran, m. and wieczorek, w., 2009. Ceramic-in-polymer vs. polymer-in-ceramic polymeric electrolytes – a novel approach. Journal of Power Sources, 194(1), 66–72. tarascon, j. and armand, m., 2001. Issues and challenges facing rechargeable lithium batteries. Nature, 414(6861), 359–367. tarascon, j., gozdz, a.s., schmutz, c., shokoohi, f. and warren, p.c., 1996. Performance of Bellcore’s plastic rechargeable Li-ion batteries. Solid State Ionics, 86–88(Part 1), 49–54.
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vaia, r.a., jandt, k.d., kramer, e.j. and giannelis, e.p., 1995. Kinetics of polymer melt intercalation. Macromolecules, 28(24), 8080–8085. van eijck, l., best, a.s., stride, j. and kearley, g.j., 2005. Softening of the potentialenergy surface in polymer electrolytes on the addition of nanoparticles. Chemical Physics, 317(2–3), 282–288. van zon, a. and de leeuw, s.w., 2001. A Rouse model for polymer electrolytes. Electrochimica Acta, 46(10–11), 1539–1544. villarreal, i., morales, e. and acosta, j.l., 1999. Ionic conductivity and spectroscopic characterisation of γ-LiAlO2-filled polymer electrolytes. Angewandte Makromolekulare Chemie, 266, 24–29. volel, m., armand, m. and gorecki, w., 2004. Influence of sample history on the morphology and transport properties of PEO–lithium salt complexes. Macromolecules, 37(22), 8373–8380. walkowiak, m., zalewska, a., jesionowski, t., waszak, d. and czajka, b., 2006. Effect of chemically modified silicas on the properties of hybrid gel electrolyte for Li-ion batteries. Journal of Power Sources, 159(1), 449–453. walkowiak, m., zalewska, a., jesionowski, t. and pokora, m., 2007. Stability of poly(vinylidene fluoride-co-hexafluoropropylene)-based composite gel electrolytes with functionalized silicas. Journal of Power Sources, 173(2), 721–728. walls, h.j., riley, m.w., fedkiw, p.s., spontak, r.j., baker, g.l. and khan, s.a., 2003. Composite electrolytes from self-assembled colloidal networks. Electrochimica Acta, 48(14–16), 2071–2077. west, k., zachau-christiansen, b., jacobsen, t. and skaarup, s., 1989. Solid-state sodium cells – an alternative to lithium cells? Journal of Power Sources, 26(3–4), 341–345. weston, j.e. and steele, b.c.h., 1982. Effects of inert fillers on the mechanical and electrochemical properties of lithium salt–poly(ethylene oxide) polymer electrolytes. Solid State Ionics, 7(1), 75–79. wieczorek, w., 1995. Composite Polyether Based Electrolytes. Warszawa: Officyna Wydawnicza Politechniki Warszawskiej. wieczorek, w., płocharski, j., przyłuski, j., głowinkowski, s. and pajak, z., 1988. Impedance spectroscopy and phase structure of PEONaI complexes. Solid State Ionics, 28–30(Part 2), 1014–1017. wieczorek, w., such, k., wycis´lik, h. and płocharski, j., 1989. Modifications of crystalline structure of peo polymer electrolytes with ceramic additives. Solid State Ionics, 36(3–4), 255–257. wieczorek, w., such, k., plocharski, j. and przyluski, j., 1990. Proc. Second Int. Symposium on Polymer Electrolytes, Scrosati, B. (ed), Elsevier, 339. wieczorek, w., florjan´czyk, z. and stevens, j.r., 1995. Composite polyether based solid electrolytes. Electrochimica Acta, 40(13–14), 2251–2258. wieczorek, w., chung, s.h. and stevens, j.r., 1996a. Ionic conduction in composite polymer electrolytes at subambient temperatures. Journal of Polymer Science, Part B: Polymer Physics, 34(17), 2911–2917. wieczorek, w., stevens, j.r. and florjan´czyk, z., 1996b. Composite polyether based solid electrolytes. The Lewis acid–base approach. Solid State Ionics, 85(1–4), 67–72. wieczorek, w., zalewska, a., raducha, d., florjan´czyk, z., stevens, j.r., ferry, a. and jacobsson, p., 1996c. Polyether, poly(N,N-dimethylacrylamide), and LiClO4 composite polymeric electrolytes. Macromolecules, 29(1), 143–155.
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wieczorek, w., zalewska, a., siekierski, m. and przyluski, j., 1996d. Modelling the a.c. conductivity behaviour of composite polymeric electrolytes by the effective medium theory. Solid State Ionics, 86–88, 357–362. wieczorek, w., lipka, p., zukowska, g. and wycis´lik, h., 1998a. Ionic interactions in polymeric electrolytes based on low molecular weight poly(ethylene glycol)s. Journal of Physical Chemistry B, 102(36), 6968–6974. wieczorek, w., zalewska, a., raducha, d., florjan´czyk, z. and stevens, j.r., 1998b. Composite polyether electrolytes with lewis acid type additives. Journal of Physical Chemistry B, 102(2), 352–360. wright, p.v., 1975. Electrical conductivity in ionic complexes of poly(ethylene oxide). British Polymer Journal, 7, 319–327. wright, p.v., 1998. Polymer electrolytes – the early days. Electrochimica Acta, 43(10– 11), 1137–1143. yang, c., kim, h., na, b., kum, k. and cho, b.w., 2006. Gel-type polymer electrolytes with different types of ceramic fillers and lithium salts for lithium-ion polymer batteries. Journal of Power Sources, 156(2), 574–580. zalewska, a., wieczorek, w. and stevens, j.r., 1996. Composite polymeric electrolytes from the PEO–PAAM–NH4SCN system. Journal of Physical Chemistry, 100(27), 11382–11388.
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3 Polymer electrolytes based on natural polymers A. PAW L I C K A and J. P. D O N O S O, Universidade de São Paulo, Brazil
Abstract: The development of new materials that can be applied as solid electrolytes has led to the creation of modern systems of energy generation and storage. Among different poly(ethylene oxide)-based electrolytes, natural polymers, such as hydroxyethylcellulose, hydroxypropylcellulose or carboxymethylcellulose, starch and chitosan or proteins like gelatin, are also proposed. After salt or acid addition, the transparent membranes show ionic conductivity values of 10−3 S/cm and can be applied with good stability results to electrochromic devices. The present review shows the principal results of ionic conductivity characterization of polymer electrolytes based on natural polymers with emphasis on solid state nuclear magnetic resonance (NMR) measurements. Key words: polymer electrolytes, natural polymers, ionic conductivity, RMN.
3.1
Introduction
Solid polymer electrolytes (SPEs) have been known and studied for over 30 years since their discovery by Peter V. Wright1 and major focus on technological potential observed by Michel Armand et al.2 The SPEs are solid solutions of monovalent (LiClO4, NaClO4, LiCF3SO3, LiBF4, etc.) or divalent (Ca(ClO4)2, MgCl2, Mg(CF3SO3)2) metal salts in polymers, of which the pioneer, first and most studied was and still is the poly(ethylene oxide) (PEO). Although the most referenced papers are those by Wright and Armand,1,2 the beginning of this work was described by Wright and co-workers in 19733 who mentioned the first temperature sensitive ionic conductivity results of the PEO–alkali metals complex. In the same year Armand4 reported ‘but it is realistic to expect that in the near future a whole set of electrolytes will be available for either Li, Na or K. Especially thin film polymers will be suitable for an all solid state system, as a good contact is easily achieved with soft materials . . .’. The enthusiasm was great because, although several solid ionic conducting systems, even with high conductivities were already known, they always presented some kind of problem, as in the case of glasses with Li which during the process of 95 © Woodhead Publishing Limited, 2010
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intercalation/deintercalation (extraction) had changes in the volume of 4–10% leading to interruptions of electrode–electrolyte contact.5 The substitution of ceramic systems by PEO-based systems made it possible to solve the interface problems and start new scientific studies of solid electrolyte directions as summarized by Armand in his paper published in 1994.5 However, despite the great enthusiasm, the PEO-based systems presented some problems, mainly because of their crystallinity and consequently low melting point, which was around 65 °C (depending on the molecular weight). This limited the development of many applications; however, thermoelectrochromic devices, like NiO/(PEO)8LiClO4/LixWO3 are still possible.6 Despite this crystallinity, most studies have continued employing the PEO in the formulation of polymer electrolytes since its ion conduction properties are better than those of poly(methylene oxide), poly(trimethylene oxide) or poly(propylene oxide).6 However, in an attempt to optimize the electrical properties and the mechanical and electrochemical stability of PEO-based systems, since 1981, new systems, generally using the PEO in the formulation, obtained by crosslinking, copolymerizations and plasticization processes have been investigated.6 Plenty of progress has been achieved in the increase in ionic conductivity, although electrochemical stability problems still remain. The SPEs have been widely studied as they are a promising alternative to replace inorganic electrolytes and liquid crystals used in batteries, sensors and various electrochemical devices, among others. As mentioned above, most systems described employ the chains of polymer that can be modified in order to decrease its crystallinity. The modification can also change the glass transition temperature (Tg) of these systems, which implies changes in the mobility of the polymer chains. The decrease in the Tg generally increases the mobility of macromolecules and consequently can improve the ionic conductivity of the material. Among various investigated systems are blends and composites7–11 with PEO, but also ionic conducting materials based on natural polymers, such as cellulose derivatives,12,13 starch,14 chitosan,15 gelatin16 and natural rubber.17 The interest in natural polymers called biopolymers is large due to the depletion of sources of finite raw materials, which is the case of oil and natural gas as well as the contribution to the reduction in the emission of industrial gases and generation of greenhouse gases. Furthermore, the use of biopolymers reduces the environmental impact due to the disposal and storage and long-term degradation of synthetic polymers since they are biodegradable and are in the organic carbon cycle providing nutrients to the soil through the decomposition process. Studies of the use of biopolymers also take into account the low production cost due to the great variety and low prices of raw materials,18 and also opportunities for the increase their use in workplaces in agriculture and the chemical industry. As an example the use of natural fibers in the preparation
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of various composites may decrease by at least 20 MJ of energy per kilogram of polymer, avoid 1 kg of CO2 per kilogram of polymer and reduce other environmental impacts by at least 20%. Concerning the development of new materials based on polysaccharides and proteins, considerable attention has been drawn in recent years. The change in the physical and chemical properties of natural polymers and improvement of their functional characteristics can be performed by means of chemical reactions such as etherification, esterification, grafting and crosslinking reactions or physical processes, like plasticization. The derivatization reactions involve the exchange or introduction of functional groups into macromolecules, leading to change in the gelatinization and retrogradation properties of these materials. As an example, the modified starch is obtained by introducing, for example, hydroxypropyl groups resulting in inter- and intramolecular hydrogen bonding interruption and, consequently, an increase in the amorphous phase.19 The preparation of different blends based on natural polymers, like polysaccharides, such as cellulose, chitosan, starch, etc. and its derivatives and gelatin, i.e. protein,14,17 allows for their processing in the film form with good mechanical properties, as well as adhesion to glass and metal surfaces. Natural polymer-based plastics can be also obtained through plasticization15,16,20,21 or grafting22 processes. The polymer electrolytes obtained by the plasticization process with low molecular weight substances increase the values of the ionic conductivity of the formed blends. Although the mechanism is not yet well known, there is an interaction of ions with mobile specific sites along the polymer chain. Also the fact that a plasticizer separates the polymer chain enhances the assumptions about their important role in the increase of the ionic conductivity values.23 Glycerin and water are widely used as plasticizers for starch, reacting as polymer chain crosslinking and separating agents. These interactions may be the van der Waals forces and hydrogen bonds formed due to polymer– polymer and polymer–plasticizer interactions.20 To obtain ionic conducting gelatin-based membrane, formaldehyde can be used to crosslink the chains and increase the stability, and glycerin as plasticizer to improve the ionic conductivity of the membranes.16
3.2
Grafted natural polymer-based solid polymer electrolytes
3.2.1 Cellulose derivatives In order to optimize the elaboration of thin devices, the electrochemical stability, conductive properties and film-forming properties of the SPEs are studied. In this field, cellulose and its derivatives are very attractive due to
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Polymer electrolytes OH O CH2OH O
O
O HO
O
HO
O CH2OH HO
O
CH
CH
3
CH2OH O
O
OC
3
m NC O
N
O
Jeffamine diisocyanate
O
HEC
OH CH3 NH m C O O
CH3 O
HOH2C HO HO
C NH O O
O H2C O
O CH2OH O
CH3
O
O CH2OH
O OH
OH OH
HO
O HO O CH3 C O NH m
O
O
O
O O CH2 OH
O HO
CH2OH O
O HO
O HEC-PEO network O OH
OCN
3.1 Hydroxyethylcellulose (HEC)–PEO network formation by grafting with Jeffamine diisocyanate.
their good film-forming, mechanical properties and rich variety in nature.24 In 1994 Le Nest et al.25 reported the possibility of obtaining solid polymeric electrolytes of hydroxyethylcellulose (HEC) and hydroxypropylcellulose (HPC) networks by grafting these polymers with diisocyanates resulting in urethane bond formation (Fig. 3.1). In 1995 Schonenberger et al.24 proposed a new kind of solid polymeric electrolytes based on HEC grafted with PEO through the condensation reaction of hydroxyl (OH) groups of HEC with PEO oligomer bifunctional isocyanates. The work makes it possible to generate a great variety of networks depending on the polymer matrix structure, its molecular weight and type of isocyanate used. Regiani et al.12,13,26 obtained and studied samples of transparent films of HEC grafted with poly(ethylene oxides) diisocyanates (Jeffamine diisocianate) shown in Fig. 3.1. After the addition of lithium salt (LiClO4) to these networks they produce solid polymeric electrolytes with ionic conductivity values of about 2.08 × 10−5 S/cm at 40 °C and 8.8 × 10−4 S/cm at 60 °C.27 Other preliminary tests concerning starch were published by Singh et al.19 and van Soest and Vliegenthart21 reporting the good mechanical and electrical properties while improving the possibility of casting films.
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99
–2 –3
log (σ S/cm)
–4 –5 –6 –7 –8 –9 –10 2.4
2.6
2.8 3.0 103/T (1/K)
3.2
3.4
3.2 Log of ionic conductivity values as a function of temperature for four HEC-based networks (䊏) HEC–PEO, (䊉) HEC–PPO, (䉱) HEC–MPPO and (夹) HEC–PhNCO.
The series of different film samples of oligo-ether isocyanates grafted HEC were analyzed by Regiani et al.28 who observed that the conductivity in these films is independent of the cellulose derivative structural parameters (molar substitution, MS, degree of substitution, DS, and degree of polymerization, DP), but it is related to the type of isocyanate used in the graft reaction. The HEC samples were grafted with monoisocyanate of poly(propylene oxide) (MPPO) and different diisocyanates of poly(ethylene oxide) (DPEO) and poly(propylene oxide) (DPPO) or phenyl isocyanate (PhNCO) in the presence of LiClO4. The samples of grafted polysaccharide derivatives showed Tg values of −40 °C for HEC/MPPO, −33 °C for HEC/ DPEO and −10 °C for HEC/DPPO. From these analyses it was observed that the Tg values depend on the type of isocyanate used to graft and crosslink HEC. Lower Tg values were observed in the grafted samples where higher values for crosslinked samples were found. The ionic conductivities of all samples were about 10−7–10−6 S/cm at room temperature and 10−4 S/ cm at 80 °C, as shown in Fig. 3.2, and it was possible to observe that the mechanism ruling the ionic conduction can be explained by the free-volume theory represented by a Williams–Landel–Ferry (WLF) equation, typical for amorphous electrolytes. Recently Chelmecki et al.29 reported the ionic conductivity results of HPC–PEO networks containing LiN(CF3SO2)2. These networks were obtained by thermal crosslinking of HPC with tri(ethylene glycol) and toluene diisocyanate through the urethane bonds, as shown in Fig. 3.1. The ionic conductivities of these samples were in the range of 10−4 S/cm. The
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Polymer electrolytes H3C
H3C
OH O
O H3C CH2OH HO O O O HO
O
O
[O]
CH2OH O
O
O CH2OH HO
H3C CH2OH HO O CH2OH O O O O O O CH2OH HO HO O
O
CH3 O
HPC
CH3 OH
CH
3
3
H3 C
3
O N
O
HO HO
mN
O
H3C HO
O H2C O
OH OH H C 2 CH2
O
HO HO
O O O CH3 O
O
O
O
Jeffamine
OH H2C
O
CH3
H3C O OH
2
CH3
m NH2
O
NH
CH
CH2
CH3 O
CH
OH
CH3
O
Oxidized HPC
O O CH2OH
O O
CH3 O
OH
CH3 N m
CH3
H2N O
O
CH3
OH HPC-PEO network
CH3
3.3 Hydroxypropylcellulose (HPC)–PEO network formation by grafting with Jeffamine.
possibility of using of these samples in lithium-ion batteries, which can be rechargeable with constant DC voltage up to the open circuit voltage of 3.65 V was also shown. A different grafting reaction was performed by Machado et al.30 using the principle of Schiff base, i.e. a reaction between primary amine and ketone groups through imine bonds, as shown in Fig. 3.3. This reaction can possibly also result in a wide variety of networks, which can be used as solid polymeric electrolytes when associated with an appropriate salt. To achieve this route it was necessary to convert HPC secondary hydroxyl into ketone groups under mild experimental conditions using 1-chlorobenzotriazol.31,32 The networks were obtained by grafting with Jeffamines (diamines with polyether chains) as shown in Fig. 3.3. HPC-based SPEs were obtained by the introduction of LiClO4 to these networks. The ionic conductivity values as a function of temperature showed two types of conductivity models, depending on the LiClO4 concentration. For low ([O]/[Li] = 50) and high ([O]/[Li] = 6) salt quantities, an Arrhenius-
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101
log (σ/S cm–1)
–3
–4
–5 6
–6
12 –7 2.6
2.8
3.0
3.2
3.4
3.6
103/ T (K–1)
3.4 Log of ionic conductivity values as a function of temperature for HPC–PEO (䊉) [O]/[Li] = 6, (䊏) [O]/[Li] = 12 and CMC–MPPO (夹) networks formation by grafting with Jeffamine.
type dependence was found, while for intermediate salt quantities of [O]/ [Li] = 30–8 Vogel–Tamman–Fulcher (VTF)-type dependence characteristic of polysaccharide ionic conducting networks was observed (Fig. 3.4). The conductivity values of the samples containing [O]/[Li] = 6 of lithium salt increased from 5.9 × 10−6 S/cm at 23 °C to 5 × 10−4 S/cm at 84 °C. However for other samples containing [O]/[Li] = 6 of lithium salt, higher ionic conductivity values of 1.26 × 10−5 S/cm at 30 °C were also observed. Figure 3.4 also shows ionic conductivity results of MPPO-grafted carboxymethylcellulose (CMC) as the SPE samples with urethane bonds with Tg of 24 °C.28 As observed in this figure, the CMC/MPPO sample presented a conductivity value similar to the one of HEC/DPPO, i.e. 10−7 S/cm; however, it did not exhibit any film-forming characteristic, indicating its use as a solid polymer electrolyte only in the pellet form.28
3.2.2 Chitosan One of the most studied polysaccharides is chitosan, due to its specific properties, such as biocompatibility and bioactivity and also its promising potential for biomedical, pharmaceutical and industrial applications.33,34 Chitosan is the N-deacetylated derivative of chitin – the structural component of shrimps, crustaceans, also found in some plants – which should have more than 50% of amine groups. As some other polysaccharide derivatives, chitosan is biodegradable and can be easily cast into thin films. Owing to its amine groups distributed regularly along the polymeric chain, in acid solution it becomes a cationic polyelectrolyte after amine groups
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protonation.35 As an SPE, it also constitutes a polymer host for ionic movements as it can dissolve ionic salts36,37 and promote ionic conductivity. In 1998 Velazquez-Morales et al.38 described the synthesis of oxipropylated chitosan grafted and crosslinked with oligoether-based mono- and diisocyanates, as in the case of cellulosic ethers. The ionic conducting membranes were obtained by introducing lithium bis(trifluoromethyl sulfonyl) imide (LiN(CF3SO2)2 – LiTFSI) by diffusion from an acetonitrile solution, yielding values similar to those obtained for the cellulose derivative networks described above.
3.2.3 Starch Starch-like cellulose is a very interesting natural polymer owing to its rich variety in nature and because it can be obtained from many renewable sources. When cellulose is the major structural component of mature plants, starch is the major reserve material of many plant storage tissues. Starch and cellulose are polysaccharides with a very similar chemical formula and a small difference of linking between glycosidic rings. In cellulose there is β-1,4 and in starch α-1,4 glucosidic bond; starch is also a mixture of the predominantly linear amylose and highly branched amylopectin.20 This dif-
Table 3.1 Summary of grafted and crosslinked natural polymer-based ionic conducting systems
Natural polymer HEC
HEC
CMC HPC HPC Chitosan Starch
σ (S/cm) at room temperature
Grafting agent
Lithium salt
Poly(ethylene oxides) diisocyanates (Jeffamine diisocyanate) Poly(propylene oxide) monoisocyanate and different poly(ethylene oxide) and poly(propylene oxide) diisocyanates Poly(propylene oxide) monoisocyanate Tri(ethylene glycol) and toluene diisocyanate Jeffamine Oligoether-based monoand diisocyanates Toluene poly(propylene oxide) diisocyanate
LiClO4
2.08 × 10−5 (40 °C)
12,26
LiClO4
10−7–10−6
28
LiClO4
10−7
28
LiN(CF3SO2)2
10−4
29
LiClO4 LiN(CF3SO2)2
10−6–10−5 10−5 (*)
30 38
LiClO4
3.5 × 10−5
*Exact data not informed.
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Reference
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ference results in an interesting property of starch, which is its solubility in some organic solvents and hot water and, consequently, a possibility of obtaining this polymer in the film form with good mechanical properties. Few results on grafted starch have been reported. Dragunski and Pawlicka39 showed results in amylopectin-rich starch samples grafted with toluene poly(propylene oxide) diisocyanate. The characterization of these film samples revealed good mechanical properties, formation of urethane bonds, decrease of crystalline phase and Tg of −11 °C. The ionic conductivity measurements reported further40 showed a change of two orders of magnitude increasing from 3.5 × 10−5 to 1.48 × 10−3 S/cm. It was also noted that the mechanism ruling the ionic conductivity follows predominantly the VTF model, and the activation energy was about 7.02 kJ/mol, which was similar to the results obtained for the networks of HEC and HPC grafted with poly(ethylene oxide) diisocyanate.28 Table 3.1 gives a summary of grafted and crosslinked natural polymerbased ionic conducting systems.
3.3
Plasticized natural polymer-based solid polymer electrolytes
The principal obstacles of the grafting and crosslinking procedures described above are the use of humidity-sensitive substances, as in the case of diisocyanates, the long time taken, expensive reagents and special conditions of synthesis, among others. To avoid these problems, another possibility of modifying the polymeric structure and decreasing the crystallinity degree of the ionic conducting samples is by using a physical process, like plasticization. This process, widely used in plastic industry, differs from chemical modification (grafting or crosslinking) principally by its ease in application, low cost and good final results. In the domain of ionic conducting polymerbased systems, those with plasticizers also have attracted the attention of various research groups. An increase in the ionic conductivity of 1 to 2 orders of magnitude and a decrease in the Tg values by 40 °C of the plasticized polymeric samples were reported by MacFarlane et al.41 and Forsyth et al.42 in 1995. The plasticization process of natural polymers using low molecular weight substances, like glycerol, ethylene glycol, propylene carbonate, sorbitol or others, also increases the values of ionic conductivity and decreases the crystallinity of the formed blends.
3.3.1 Starch Starch, cellulose and other polysaccharides can be modified by chemical or physical processes such as grafting with isocyanates or plasticization with low molecular weight substances.20 The addition of plasticizer to starch reduces its crystallinity and Tg, which is a very important parameter to © Woodhead Publishing Limited, 2010
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obtain solid polymeric electrolytes. Water and glycerol are widely used as plasticizers for starch and act as agents of polymeric network formation through the polymer–polymer and polymer–plasticizer interactions. These interactions are van der Waals forces and strong hydrogen bonds, which lead to the increase in the toughness and strength of the starch plasticized with glycerol materials.21 In 2002 Dragunski and Pawlicka43 presented the first results of new solid polymeric electrolytes based on amylopectin-rich starch plasticized with 30 wt% of glycerol and containing different quantities of LiClO4. The transparent film samples showed low Tg values of about −12 °C and a predominantly amorphous state even after the addition of lithium salt, which was totally dissolved in the polymer matrix. The measurements of ionic conductivity as a function of temperature showed an Arrhenius-type dependence and an increase of two orders of magnitude in the temperature interval from 30 to 80 °C, i.e. from 5.05 × 10−5 S/cm at 30 °C to 7.15 × 10−3 S/cm at 80 °C for the sample with lithium salt concentration of [O]/[Li] = 8 (Fig. 3.5). The same figure shows the results of the SPE based on starch plasticized with ethylene glycol and containing lithium salt with concentration of [O]/[Li] = 12, where the ionic conductivity values increase from 7.91 × 10−5 S/cm to 3.32 × 10−4 S/cm in the same temperature interval. Investigations into the influence of humidity on the ionic conductivity values of the ionic conducting samples based on plasticized starch were reported by Rodrigues and Pawlicka.44 The authors of this study showed that the increase in the humidity of the samples promotes an increase in the ionic conductivity; the sample with salt concentration of [O]/[Li] = 11
–2
[O]/[Li] 12 8 6 starch-EG
log (σ/S cm–1)
–3
–4
–5
–6 2.7
2.8
2.9
3.0
3.1 3.2 103 K/T
3.3
3.4
3.5
3.5 Log of ionic conductivity measurements as a function of the inverse of temperature for the samples of amylopectin plasticized with glycerol (30 wt%) and containing [O]/[Li] = 6 (*), 8 (䉲) and 12 (䊏) and plasticized with 30 wt% of ethylene glycol [O]/[Li] = 12 (䊏).
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changes from 3.88 × 10−5 S/cm at 22% of humidity to 9.45 × 10−4 S/cm at 55% of humidity. They also observed that the ionic conductivity values of these samples were very stable for several hours, principally at low humidity values and low salt concentration ([O]/[Li] = 14). At higher relative humidity values and higher salt concentration ([O]/[Li] = 11), the conductivity values were less stable, indicating the influence of water on the ionic conductivity of the samples. Following the supposition of Nogueira et al.,45 this ionic conductivity increase can be attributed to the presence of water, which may have given rise to a new route of ionic transport facilitated by the higher diffusion coefficient or promotion of plasticization effect in the polymer chain and, consequently, its better mobility. It can also be due to the proton conduction, as concluded in the humidity effect on the ionic conductivity of PEO-based SPEs study reported by Tanzella et al.46 Aiming to verify the influence of starch source and modification on polymer electrolytes’ properties, different SPEs based on native and modified corn and cassava starches were obtained and characterized.47 In this study it was stated that the ionic conductivity values of all analyzed samples varied in the range from 10−6 to 10−4 S/cm, and the best ionic conductivity values were obtained for the samples of oxidized and acetylated starches plasticized with 30 wt% of glycerol and containing 8 and 10 [O]/[Li] lithium salt concentrations. The values were 1.75 × 10−4 S/cm at 32 °C and 1.15 × 10−3 S/cm at 85 °C for the oxidized starch-based SPEs, and 4.54 × 10−4 S/cm at room temperature and 2.4 × 10−3 S/cm at 80 °C for acetylated starch-based SPE samples. The poor ionic conductivity values of 1.62·10−6 S/cm at room temperature were obtained for the cationic starchbased SPEs.
3.3.2 Agar Another possible candidate as a matrix for natural polymer-based electrolytes is agar, which like starch is a heterogeneous mixture of two polysaccharides: agaropectin and agarose. Although both polymers share the same galactose-based backbone, agaropectin is heavily modified with acidic sidegroups, such as sulfate and a pyruvate, while agarose has a neutral charge. The purpose of this study was the obtaining of protonic conducting matrix using acetic acid in polymer electrolyte formulation.48 The membrane samples showed 70–92% of transparency in the visible range depending on the acetic acid concentration. It was also found that the acetic acid amount greatly influences the ionic conduction. The best ionic conductivity values of 1.1 × 10−4 S/cm at room temperature and 9.6 × 10−4 S/cm at 80 °C were obtained with this type of sample. This study also verified the influence of relative humidity on ionic conductivity values confirming a decrease to 2 × 10−5 S/cm in an atmosphere with a 33% controlled humidity. As in other
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plasticized samples studied, the ionic conductivity results as a function of temperature exhibited an Arrhenius behavior and the best values of activation energy were 33.5 kJ/mol.
3.3.3 Cellulose derivatives Cellulose ether derivatives are very important substances due to their good viscosimetric, emulsifying, stabilizing, dispersing and agglutinant properties, as well their high solubility in some organic solvents and water and ability to form films with desirable mechanical properties.49 These characteristics make cellulose ether derivatives appropriate for applications in ink, paper, food, pharmaceutical, cosmetic, ceramic, textile and agricultural products. Hydroxypropylcellulose ester In 2003 Yue et al.50 reported on gel and film electrolytes based on commercial cellulose-based product modified by a plasticization and/or crosslinking diisocyanate reaction. They studied samples composed of a hydroxypropylcellulose ester with poly(oxyethylene) side chains plasticized with propylene carbonate (PC), or a mixture of propylene carbonate with ethylene carbonate (EC) and containing lithium trifluoromethane sulfonate salt (LiCF3SO3). The ionic conductivities of gels were in the order of 10−3 S/cm for the samples plasticized with propylene carbonate contents above 50 wt% and containing 10 wt% of LiCF3SO3 and exhibited an Arrhenius behavior with activation energies ranging from 16 to 34 kJ/mol. Gels containing 70 wt% PC, or mixtures of PC and EC, were crosslinked with 1,6-diisocyanatehexane to form elastomeric non-porous films (for [NCO]/ [OH] = 2 : 1) with good mechanical properties and room temperature conductivities also of about 10−3 S/cm. The conductivities and mechanical strength of these films increased with increasing ethylene carbonate content. An Arrhenius dependence of the log of ionic conductivity on the inverse of temperature is also observed in these films, with activation energies for ion transport of 16 kJ/mol. Hydroxyethylcellulose Solid polymeric electrolytes based on HEC plasticized with different quantities of glycerol and containing LiCF3SO3 salt were reported by Machado et al.51 The samples were prepared in the transparent film form, where the transparency depended on the plasticizer quantity, with very good adhesion properties and best ionic conductivity values of 1.07 × 10−5 S/cm at 30 °C and 1.06 × 10−4 S/cm at 83 °C for the samples plasticized with 48 wt% of glycerol and containing [O]/[Li] = 6 of LiCF3SO3. It was demonstrated that
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the samples are predominantly amorphous and the Tg of the synthesized materials decrease with the increase in the quantity of glycerol from −50 °C for the sample plasticized with 12 wt% of glycerol to −83 °C for the sample plasticized with 48 wt% of glycerol. A subsequent study was reported by Machado et al.,52 concerning the influence of plasticizer contents on the ionic conductivity of HEC-LiClO4 based samples. It was verified that the best conductivity values were obtained for samples with higher quantities of glycerol when compared with plasticized starch samples,43 i.e. for the samples with more than 48 wt% of glycerol. The authors compared the ionic conductivity values at room temperature (22 °C) of the samples with and without 48 wt% of plasticizer, which decreased from 6.2 × 10−5 to 5.9 × 10−10 S/cm, respectively. It can also be stated that the temperature dependence on the ionic conductivity of both systems obeys the Arrhenius law, giving the values of activation energies (Ea) of 41 kJ/mol for the sample with 48 wt% of glycerol and 86 kJ/ mol for the sample without glycerol. From these measurements one can state that the presence of plasticizer decreases the Ea of the SPE sample and promotes a considerable increase in its ionic conductivity values. Among many different samples the best ionic conductivity results were obtained with the sample plasticized with 48 wt% of glycerol and containing [O]/[Li] = 8. Values of 9.6 × 10–5 S/cm were found, confirming other studies, in which a maximum of ionic conductivity values as a function of lithium salt concentration between [O]/[Li] = 643,51 and [O]/[Li] = 11 was also observed.53
3.3.4 Chitosan The first report on the electrical properties of polymer electrolytes based on chitosan was made by Morni et al.54 who showed ionic conductivity of 2.6 × 10−5 S/cm for the samples containing AgNO3. Osman et al.15 and Arof et al.55 presented the results of EC plasticized chitosan acetate, which exhibits the conductivity values of about 10−5 S/cm. Other samples studied in the same laboratory by Yahya and Arof37 revealed that chitosan plasticized with 50 wt% EC and the samples plasticized with 10 wt% of oleic acid showed ionic conductivities of 10−5 S/cm. The transparent membranes of chitosan-based polymer electrolytes with good ionic conductivity properties were obtained by the casting technique from a solution containing chitosan, hydrochloric acid and plasticizer such as glycerol, ethylene glycol and sorbitol.56 The best ionic conductivity values of 9.5 × 10−4 S/cm at room temperature and 2.5 × 10−3 S/cm at 80 °C were obtained for the sample containing 59 wt% of glycerol and an equimolar amount of HCl in respect to NH2 groups in chitosan. The temperature dependence of the ionic conductivity exhibits an Arrhenius behavior with
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activation energy of 16.6 kJ/mol. The thermal analysis indicates that both Tg (−87 °C) and crystallinity are low for this electrolyte. However after the addition of 13 wt% of LiCF3SO3 to the sample formulation, a decrease in the ionic conductivity values was observed, i.e. 2.2 × 10−5 S/cm at room temperature and 4 × 10−4 S/cm at 80 °C.
3.3.5 Rubber Another natural polymer-based ionic conducting plasticized system is the modified natural rubber presented by Idris et al.57 The authors obtained and characterized the samples of 25% and 50% epoxidized natural rubber (NR) and poly(methyl methacrylate) (PMMA) grafted NR unplasticized and plasticized with PC and EC and containing LiCF3SO3. They showed that the unplasticized samples exhibit ionic conductivity values of about 10−6– 10−5 S/cm, which increase to up to 10−4–10−3 S/cm for plasticized samples. They also supposed that epoxy groups of natural rubbers dissociate the salt, similarly to the PEO, and mentioned the very good adherence property to the steel electrodes of the samples, which is interesting for eventual possible application. In their following work Glasse et al.58 showed the results of systems containing epoxidized NR, PEO and salt. These samples were modified by the addition of plasticizer and the results were comparable with those of the previous work, i.e. 10−5 S/cm for unplasticized and 10−4 S/cm for plasticized samples. Kumutha and Alias59 performed a Fourier transform infrared (FTIR) study of the samples of NR grafted with 30% of PMMA plasticized with EC and containing lithium trifluoromethanesulfonate (LiCF3SO3). The aim of this study was to investigate the interactions between the polymer and salt and also the influence of the plasticizer (EC), on the salted polymer electrolyte. From this study the authors observed the complex formations in NR–LiCF3SO3 and EC–LiCF3SO3 systems, but no interaction between NR–EC. It was also concluded that the plasticizer molecules had penetrated the salted polymer matrix in NR–LiCF3SO3–EC complex.
3.3.6 Gelatin The ability of proteins to form films or membranes has been known for many years.60 Gelatin was also one of the first materials to be used in the production of movies. As other natural polymers, this protein is very promising in the area of development of new materials, since it is abundant and biodegradable, has low cost and good film-forming properties, is not toxic and forms transparent solutions with high viscosity.61 Gelatin films are generally obtained by solubilization, heat and dehydration of collagen, which leads to a distortion of a partial triple helix of this macromolecule. However
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Table 3.2 Summary of plasticized natural polymer-based ionic conducting systems
Natural polymer
Plasticizer
σ (S/cm) Ion conducting room agent temperature
Starch Starch Oxidized starch Acetylated starch Cationic starch Agar HPC ester–PEO HPC ester–PEO crosslinked with 1,6-diisocyanatohexane HEC HEC Chitosan acetate Chitosan Chitosan Chitosan Rubber–PMMA Rubber–PEO Rubber–PMMA Gelatin Gelatin Gelatin
Glycerol EC Glycerol Glycerol Glycerol Glycerol PC PC with EC
LiClO4 LiClO4 LiClO4 LiClO4 LiClO4 Acetic acid LiCF3SO3 LiCF3SO3
5.05 7.9 1.75 4.54 1.62 1.1 10−3 10−3
Glycerol Glycerol EC EC, oleic acid Glycerol Glycerol PC, EC PC, EC EC Glycerol Glycerol Glycerol
LiCF3SO3 LiClO4
1.07 × 10−5 9.6 × 10−5 10−5 10−5 9.5 × 10−4 2.2 × 10−5 10−4–10−3 10−4
HCl HCl + LiCF3SO3 LiCF3SO3 LiCF3SO3 LiCF3SO3 CH3COOH LiClO4 LiBF4
× × × × × ×
10−5 10−5 10−4 10−4 10−6 10−4
4.5 × 10−5 1.5 × 10−5 10−5
Ref. 43 47 47 47 48 50 50
51 52 15,55 37 56 56 57 58 59 64 65 66
Note: EC, ethylene carbonate; PC, propylene carbonate.
to improve its functional properties, crosslinking agents such as formaldehyde, glutaraldehyde or glioxal are used.62,63 Three different polymer electrolytes based on crosslinked gelatin with formaldehyde, and plasticized with glycerol and containing acetic acid,64 LiClO465 and LiBF466 were obtained and characterized by Vieira et al. It was found that the amount of acetic acid in proton-conducting samples influences the proton conduction and high ionic conduction values of 4.5 × 10−5 S/cm at room temperature and 3.6 × 10−4 S/cm at 80 °C were obtained with 26.3 wt% acetic acid content. Also in the case of the samples with LiClO4, the amount of lithium influences the electrolyte conductivity with a maximum for 7.9 wt% of lithium salt. However, as in the previous study, for this optimum salt concentration the ionic conductivity values increase from 1.5 × 10−5 S/cm at room temperature to 4.9 × 10−4 S/cm at 80 °C. A more recent study using factorial design type 22 with two variables, glycerol and LiBF4, of polymer electrolytes based on gelatin and LiBF4
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revealed that the effect of plasticizer on the ionic conductivity results is much more important than the effect of lithium salt or the effect of the interaction of both variables. Also in this case the ionic conductivity values at room temperature were in the range of 10−5 S/cm. The ionic conductivity results as a function of temperature of all gelatin-based SPE samples exhibited an Arrhenius behavior. These samples applied to 6 × 8 cm2 electrochemical devices showed very good performances, i.e. 25 000 switching cycles (±2.0 V/90 s) and a transmission change at 550 nm between 60% (bleached state) and 40% (colored state) corresponding to a change in the optical density (ΔOD = 0.15) with a coloration efficiency increasing from 10 cm2/C (initial cycle) to 23 cm2/C (25 000th cycle).67 A summary of the plasticized natural polymer-based polymer electrolytes is presented in Table 3.2.
3.4
Other natural polymer-based systems
The electrical properties of polymer electrolytes based on chitosan/PEO blends with LiTFSI salt were described by Idris et al.68 Conductivities between 10−5 and 10−4 S/cm were reported for samples based on swelling chitosan membranes and ammonium salts (NH4NO3 and NH4CF3SO3), which in water promote the protonation of chitosan amino groups, leading to protonic conductivity.69,70 Ionic conducting composites of chitosan, poly(aminopropyltriethoxysil ane), poly(ethylene oxide) (CHI/pAPS/PEO) and a fixed amount of lithium salt were studied by Fuentes et al.71 The ternary composition Lix(CHI)1(PEO)2(pAPS)1.2 reached a conductivity value of 1.7 × 10−5 S/cm at room temperature. It was also stated that this very transparent and good conducting structure was constituted essentially by lithium ion coordinated by APS residues intercalated in a layered CHI/PEO matrix. The ionic conductivity in the wet state of phosphorylated chitosan membranes prepared from the reaction of orthophosphoric acid and urea on the surface of chitosan membranes in N,N-dimethylformamide was investigated by Wan et al.72 The authors observed that similarly to unmodified chitosan membranes phosphorylated chitosan membranes are hardly conductive in their dry states with conductivities between 10−9 and 10−10 S/cm. The entire conduction process occurs after the water incorporation increases the ionic conductivity values up to 10−5 and 10−4 S/cm depending on the phosphorus content in the sample. The best result of 1.2 × 10−3 S/cm was obtained with the sample containing 87.31 mg/m2 of phosphorus content. They also observed that the increase in the phosphorus content promotes a decrease in the crystallinity of phosphorylated chitosan membranes, an increase in the swelling index and not a significant loss of their tensile strength and thermal stability in comparison with the unmodified chitosan membranes.
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Table 3.3 Summary of the different natural polymer-based ionic conducting systems Natural polymer-based SPE system Chitosan/PEO–LiN(CF3SO2)2 Chitosan, poly(aminopropyltriethoxysilane), poly(ethylene oxide) Phosphorylated chitosan–water di-o-butyrylchitosan–water Cyanoethylated cellulose, multifunctional poly(oxyethylene) methacrylates Chitosan/polyaniline Acacia gum and cationic HCl doped polyaniline
σ (S/cm) −5
−4
Ref.
10 –10 1.7 × 10−5
68 71
10−5–10−4 10−4–10−3 2.7 × 10−3
72 73 74
10−7–10−4
75 76
Another study published by Wan et al.73 reported on the di-o-butyrylchitosan, prepared by reacting chitosan with butyric acid anhydride in the presence of perchloric acid as a catalyst. As in the previous study, the authors observed a decrease in the crystallinity of the samples due to the modification reaction and that the swelling indices of modified membranes were increased significantly in direct proportion to the degree of substitution. The thermogravimetric analysis indicated that the modified membranes exhibited a slightly increased thermal stability compared to the unmodified membrane. The ionic conductivity of di-o-butyrylchitosan membranes in the dry state exhibited ionic conductivities between 10−10 and 10−11 S/cm, which increased up to 10−3–10−4 S/cm for 25% of substitution after hydration. Gel polymer electrolytes consisting of high polarity cyanoethylated cellulose derivatives and several kinds of multifunctional poly(oxyethylene) methacrylates and ionic conductivity of 2.7 × 10−3 S/cm at 25 °C were described and applied to a lithium-ion gel polymer battery by Sato et al.74 A large (2500 mA h) cell showed good discharge performance even in a low temperature discharge and improved safety characteristics as judged by a nail penetration test. Furthermore, this battery system exhibited a unique phenomenon of preventing overcharging. The new design of this lithium-ion gel polymer battery could be promising for large batteries that must be inherently safe, such as batteries for mobile applications. Chitosan/polyaniline (PANI) semi-interpenetrating network (semi-IPN) polymers were prepared and studied by Kim et al.75 The authors observed that the conductivity of the semi-IPNs increases from 10−7 to 10−4 S/cm with increasing PANI contents, adjusted to pH < 1, forming a blended structure. They stated that the increase in the electrical conductivity caused by the interaction of the components also reflects the charge transfer and is associated with the acidic doping of PANI.
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Electrically conductive zwitterion hybrid materials were obtained using anionic acacia gum (AG) and cationic HCl doped PANI by means of radical copolymerization method by Tiwari et al.76 This AG–PANI (acid doped) zwitterion material showed a semiconducting behavior and followed 1D (nearest neighbor) Mott’s variable range hopping (VRH) conduction mechanism in the temperature range of 83–323 K. Table 3.3 shows a summary of the other natural polymer-based polymer electrolytes presented.
3.5
Magnetic resonance spectroscopy of polymer electrolytes obtained from natural polymers
Nuclear magnetic resonance is a well-known experimental technique to study the ionic and the polymer dynamics in SPE. As an element-selective method that is sensitive to local interactions, measurements of NMR line shapes and spin-lattice relaxation times as a function of temperature provide important information on the atomic and molecular motions that modulate these interactions. Numerous reports have been published over the last three decades concerning the NMR of polymer electrolytes consisting of lithium salts dissolved in PEO. Reported NMR studies have been recently reviewed by Böhmer et al.77 In these systems, the lithium ions are coordinated to oxygen sites on the PEO so that their dynamics is strongly correlated with the polymer segmental motion. Early investigations used the NMR technique to determine the phase behavior of polyether electrolytes, i.e. the fraction of nuclei in crystalline and amorphous phases of the SPE.6,78 In the last two decades, NMR techniques have been extensively used for the study of lithium mobility as well as the segmental motion of the polymer chains of polymer electrolytes, plasticized SPE, polymer gel electrolytes and nanocomposite gel electrolytes.79–86 Although there exists a considerable number of NMR studies concerned with PEO-based SPE, few studies have been dedicated to the SPE obtained from natural polymers. Regiani et al.27 and Tambelli et al.13 reported the differential scanning calorimetry (DSC), X-ray diffraction (XRD), complex impedance spectroscopy and NMR characterization of HEC : polyether films. The experiments were performed in networks of HEC grafted with poly(ethylene oxide) diisocyanate (DPEO) or poly(propylene oxide) diisocyanate (DPPO) complexed with lithium salts. The NMR experiments exhibit the qualitative features associated with the lithium ion and polymer chain motilities in polymer electrolytes, namely the presence of well-defined 7 Li and 1H spin-lattice relaxation rate maxima (T 1−1). The single absorption band in the proton spectra and the observed 1H and 7Li exponential relaxation indicate that the HEC/DPEO : LiClO4 polymer electrolytes were predominantly amorphous.13,27
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It was realized early that the incorporation of plasticizers into polymer electrolytes reduces their crystallinity and glass transition temperatures, and enhances the ionic conductivity by reducing ionic aggregation. Lopes et al.87 investigated the conductivity, 1H and 7Li NMR spectral line shapes as well as spin-lattice relaxation times of the polymer gel electrolytes formed by HEC plasticized with glycerol and containing lithium perchlorate. The 7Li NMR results indicate that the enhancement of the conductivity with the addition of the solvent was associated with a decoupling of the lithium ion dynamics from the local motion of the polymer host, with the increase in the mobility of the ionic species and the increase in the charge carrier concentration resulting from the salt dissociation after the addition of the solvent. The temperature dependence of the lithium NMR spin-lattice relaxation was interpreted assuming that there are two distinct lithium ion dynamics. Lopes et al.88 also examined the polymer gel electrolytes formed by amylopectin-rich starch plasticized with glycerol and containing lithium perchlorate by complex impedance spectroscopy and NMR. Starch-like cellulose is an interesting natural polymer due to its rich variety and abundance in nature and good solubility properties. The heteronuclear decoupling NMR experiments suggest that the cation–proton interaction in the plasticized electrolyte was smaller than that in the grafted systems. The position of the 7Li spin-lattice relaxation rate maximum indicates that lithium has a relatively high mobility in the plasticized electrolytes. Mattos et al.89 studied the effect of water absorption on the same polymer gel electrolytes and observed a shift in the position of the 7Li spin-lattice relaxation rate maximum towards lower temperatures with increasing hydration. The shift was attributed to the increase in the lithium mobility with the hydration of the polymer gel. The 7Li relaxation results of the hydrated gel electrolytes suggested that the lithium ions interacted with the water molecules and probably dissolved into a new phase. Mattos et al.90 also reported an NMR and complex impedance spectroscopy study of a gelatin-based SPE plasticized with glycerol and containing lithium perchlorate. Gelatin also constitutes a polymer host for electrolyte as it can dissolve ionic salts. Ionic conductivities of about 10−5 S/cm were obtained at room temperature for samples prepared with 8–10 wt% of lithium salt. The temperature dependence of the 7Li spin-lattice relaxation suggested that there are two dynamic processes associated with lithium diffusion in the plasticized gelatin-based electrolytes. Donoso et al.91 investigated the lithium dynamics of a series of polymer electrolytes formed by PEO, chitosan (QO), amino propil siloxane (pAPS) and lithium perchlorate by means of 7Li NMR. Chitosan has been under extensive research on account of its specific properties, such as biocompatibility, and because of its promising potential for biomedical applications.
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Two different electrolytes were studied, a binary system (PEO/QO/Li+) and a ternary one (PEO/QO/pAPS/Li+). In the binary system the NMR data were analyzed by assuming two lithium dynamics – the first associated with the motion of the lithium ion dissolved in PEO and the second associated with those complexed by the chitosan. Although the ionic dynamics of proton conducting polymer gel electrolytes have been extensively studied in the last decades,92–95 few studies have addressed the complexes formed by natural polymers and acids. Recently Mattos et al.96,97 reported conductivity and proton (1H) NMR line shapes and spin-lattice relaxation measurements in a series of conducting gel electrolytes based on a gelatin network containing acetic acid (CH3COOH) plasticized with glycerol and crosslinked with formaldehyde. The NMR results showed that the proton mobility is dependent on the acetic acid content in the plasticized polymer gel electrolytes. In order to investigate the coordination environment of the cation in this type of gel electrolyte the authors also used the electron paramagnetic resonance (EPR) technique with samples doped with copper perchlorate. EPR is a well-known sensitive spectroscopic technique for the study of the location, nature and electronic properties of Cu2+ ions. The EPR results revealed the presence of two different copper species and the electron spin echo envelope modulation (ESEEM) results showed the interaction of 1H and 14N nuclei with the Cu2+ ions.96,97 Recently an NMR study of the proton spin-lattice relaxation as a function of temperature in a series of new proton conducting gel electrolytes was performed based on agar containing acetic acid plasticized with glycerol and crosslinked with formaldehyde.98 The temperature dependence of the 1 H spin-lattice relaxation showed a well-defined rate maximum whose position depends on the acetic acid concentration. The NMR relaxation curve was similar to those observed in the gelatin-based gel electrolyte, reflecting the high proton mobility in these gel electrolytes. Among the polymer electrolytes obtained from natural polymers, Klinklai et al.99 reported ionic conductivity and 13C NMR spectroscopy of polymer electrolytes formed by the mixture of liquid epoxidized natural rubber with lithium salts. The authors discussed the correlation between the 13C spinlattice relaxation times with the solubility of the different lithium salts in the deprotonized liquid rubber. Based on the 13C NMR chemical shift and the resonance half-width and the spin-lattice relaxation times results they concluded that lithium bis(trifluoro methane sulfonyl)imide is a suitable salt for forming polymer electrolytes with deprotonized liquid natural rubber. In a different approach, Samir et al.100 used cellulosic nanocrystals (whiskers) as organic fillers in PEO–Li[(CF3SO2)2N] polymer electrolytes. The authors investigated the ionic conductivity, electrochemical stability, ionic
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Li - 155.43 MHz
(a)
(b)
–100
–50
0
50
100
Frequency (kHz)
3.6 Low temperature (183 K) 7Li NMR spectrum of the polymer electrolyte formed by hydroxyl ethyl cellulose (HEC) grafted with PEO– diisocyanate (DPEO) complexed with lithium perchlorate: (a) 7Li NMR spectra (b) 7Li–{1H} decoupled spectra.
mobility and thermomechanical behavior of the nanocomposite polymer electrolyte and concluded that the mechanical reinforcement does not affect the lithium transference number. Some experimental results obtained from NMR and EPR spectroscopies in a series of polymer electrolytes based on cellulose derivatives, chitosan, starch, agar and other natural polymers are now presented.
3.5.1 Nuclear magnetic resonance line shape analysis Figure 3.6(a) shows a typical low temperature 7Li NMR spectrum observed in polymer electrolytes obtained from natural polymers, in this case a sample formed by HEC grafted with DPEO complexed with lithium perchlorate. The 7Li spectra below the Tg are characterized by a narrow main absorption (full-width at half-maximum ΔH ~ 7 kHz) superimposed on a broader baseline (ΔH ~ 20 kHz). The NMR spectrum of a spin I = 3/2 system (as 7Li) in a powdered crystalline sample is expected to consist of a narrow central component due to the 1/2 ↔ −1/2 transition and a doublet due to the 3/2 ↔ 1/2 and −1/2 ↔ −3/2 transitions. Nevertheless, in polymer electrolytes, one usually observes a distribution of ±3/2 ↔ ±1/2 nuclear quadrupole satellite transitions below Tg, resulting in a shoulder on each side of the central ±1/2 transition. This fact is attributed to a distribution of
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electric field gradients (EFG) at the lithium site due to the heterogeneous Li nearest-neighbor configuration in the polymer electrolyte. This last interaction is parameterized in terms of the quadrupole coupling constant νq whose mean value can be estimated from the half-width of the broad baseline in the 7Li NMR absorption spectrum. The width of the central line associated with the 1/2 ↔ −1/2 transition is determined by heteronuclear magnetic dipole–dipole couplings interaction between the lithium nucleus and the surrounding nuclei. Besides the homonuclear dipole–dipole broadening of the central line, there are heteronuclear contributions from protonic species (HEC and DPEO) and coupling with the chlorine and oxygen from ClO4 species. Given the small gyromagnetic factors and/or low natural abundance of 13C, 35Cl and 17O isotopes, it was concluded that the homonuclear 7Li–7Li interaction and the heteronuclear 7Li–1H interaction are responsible for the lithium linewidth observed. The NMR decoupling technique was used to discriminate between these interactions. The 7Li–{1H} decoupling experiments performed at sufficiently low temperature (173 K) demonstrated that 70% to 95% of the sub-Tg 7Li linewidth account for the heteronuclear 7Li–1H interaction (Table 3.4). A reduction in the 7Li linewidth (from 7.2 to 1.6 kHz) is clearly visible in the central line in Fig. 3.6(b), reflecting the strong interaction between the 7Li and the protons of the HEC and DPEO. Table 3.4 gives the NMR parameters obtained from the 7Li NMR experiments in polymer electrolytes. Figure 3.7 shows the typical behavior of the 7Li central linewidth as a function of temperature for gel electrolytes obtained from natural polymers. The first one is based on amylopectin-rich starch plasticized with glycerol and containing lithium perchlorate89 and the second one formed by PEO, chitosan and lithium perchlorate.91 Above 170–200 K, the lithium mobility increases sufficiently to average dipolar interactions, producing a line narrowing. This motional narrowing is attributed to the modulation of
Table 3.4 NMR parameters obtained from the 7Li NMR experiments in polymer electrolytes. Δω is the low temperature 7Li linewidth and Δω(7Li–1H) is the residual 7 Li linewidth after proton decoupling Polymer electrolyte
Δν (7Li) kHz
Δν 7Li–{1H}
Li–H
Reference
PEO8LiClO4 HEC/DPEO/LiClO4 Amylopectin/glycerol/LiClO4 PEO/chitosan/Li Gelatin/glycerol/LiClO4
6.4 7.2 7.8 6.6 6.7
0.5 1.6 2.4 0.3 0.9
90% 78% 70% 95% 87%
80 13 88 91 90
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8
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Li 155.43 MHz
ΔH (kHz)
6 4 2 0 150
200
250
300
350
400
T (K)
3.7 Temperature dependence of the 7Li NMR central linewidth for the polymer gel electrolytes based on amylopectin-rich starch plasticized with glycerol and containing lithium perchlorate (䊉) and the electrolyte formed by PEO, chitosan and lithium perchlorate (䉱).
the 7Li–1H magnetic dipole–dipole interaction caused by the lithium motion. The criterion for motional narrowing is that the rate of fluctuation of the local lithium dipolar fields or the electric field gradients, generally described in terms of a correlation time τc, becomes much higher than the low temperature rigid lattice linewidth (about 8 kHz). The observation of a single, motionally averaged resonance in Fig. 3.7 indicates that the lithium ions in the electrolytes must be undergoing rapid diffusion with a 1/τc rate. In most cases, the motional narrowing takes place near the calorimetric Tg. This correspondence between the Tg and the onset of the NMR line narrowing has been attributed to the coupling between ion dynamics and polymer segmental motions in polymer electrolytes.77 For the electrolytes in Fig. 3.7, for example, the onset of motional narrowing occurs at about 220 K for amylopectin plasticized with glycerol electrolyte and at about 180 K for the sample of PEO : chitosan. It should be noted that in plasticized polymer gel electrolytes, the 7Li narrowing process could be decoupled from the glass transition. In these systems the lithium mobility can be associated with the mobility of the plasticizer molecules and the 7Li line transition should occur at the onset of the reorientational dynamics of the solvent molecules. From the 7Li line narrowing data it is possible, in principle, to deduce the temperature dependence of the correlation time, τc, associated with the lithium motion. Assuming a thermally activated process for the lithium motions following an Arrhenius temperature dependence for τc: τc = τ0 exp(Ea/kBT)
[3.1]
where kB is the Boltzmann constant and τ0 is a pre-exponential factor, the activation energy Ea for the line narrowing processes can be obtained.88,91
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In general, the activation energies determined from line narrowing analysis are larger than those obtained from NMR spin-lattice relaxation measurements. While the NMR line narrowing process occurs in a narrow temperature region near Tg, the 7Li spin-lattice relaxation is examined at high temperature, above Tg. Böhmer et al.77 pointed out that the assumption of the existence of a thermally activated process described by a single correlation time is not necessarily found in all polymer electrolytes, so that care should be taken when interpreting experimental results. Since the main contribution to the width of the central 7Li NMR line results from 7Li-1H dipolar couplings, the line narrowing can be a consequence of both lithium and polymer dynamics, which occur on the same timescale in the amorphous state.
3.5.2 Nuclear magnetic resonance spin-lattice relaxation Figure 3.8 displays the Arrhenius plots of the 7Li spin-lattice relaxation rates (T 1−1) for three electrolytes based on natural polymers complexed with lithium perchlorate, all measured at the Larmor frequency of 155 MHz. The
101 7
1/T1 (s–1)
Li - 155.4 MHz
100
10–1 2.5
3.0
3.5 4.0 4.5 1000/T (K–1)
5.0
3.8 Temperature dependence of the 7Li spin-lattice relaxation rates (T 1−1) for the polymer gel electrolytes based on amylopectin-rich starch plasticized with glycerol and containing LiClO4 (䊏), the electrolyte formed by PEO, chitosan and LiClO4 (䉱) and the polymer electrolyte formed by HEC grafted with DPEO complexed with LiClO4 (䊊), measured at the Larmor frequency of 155.4 MHz.
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first electrolyte is formed by HEC grafted with DPEO,13 the second is based on amylopectin-rich starch plasticized with glycerol89 and the third is formed by PEO and chitosan.91 A single relaxation rate maximum was observed in all the samples, but the relaxation curves of the last electrolyte exhibited an asymmetric shape around the rate maxima. A well-defined T1−1 maximum indicates that only one motion process of the lithium ion is taking place in the electrolyte. In contrast, an asymmetric shape of the relaxation curve suggests that there are at least two separate relaxation processes in the electrolyte. The temperature dependence of the spin-lattice relaxation rate T 1−1 is usually analyzed using the simple Bloembergen, Purcell and Pound (BPP) model, which assumes a single correlation time describing non-correlated isotropic random motions.77,80 The spin-lattice relaxation expressed in terms of the spectral density function J(ω) evaluated at the NMR Larmor frequencies ω0 and 2ω0 is 1 = C [ J (ω0 ) + J ( 2ω0 )] T1
[3.2]
where constant C depends on the particular spin interaction responsible for the relaxation. The first step in 7Li NMR spin-lattice relaxation studies is to establish the dominant relaxation mechanism. The 7Li NMR relaxation processes in lithium salt polyether complexes are governed mainly by two mechanisms: (i) the quadrupolar relaxation, due to the interactions between the quadrupole moment of the 7Li nuclei and fluctuations in the EFG produced by the charge distribution at the site of the nucleus, and (ii) the dipolar relaxation, produced by random fluctuations of the lithium homoand heteronuclear dipole–dipole interactions. The dipolar contribution to the spin-lattice relaxation rate can be calculated from the second moment of the static line-shape obtained from the 7Li spectrum. By comparing the result with the measured value of the relaxation rate maximum, the quadrupolar contribution can be estimated. This simple calculation demonstrates that fluctuations of the quadupolar interaction due to the Li+ motions are mostly responsible for the 7Li NMR spin-lattice relaxation observed in SPE.13,91 This analysis is a reasonable approximation for ion dynamics in polymer electrolytes providing that there is no distribution of correlation times (which leads to a decrease in the value of T 1−1 at the rate maximum).77 The spectral density function, which describes the density of fluctuations in the magnetic field, is given by J (ω ) =
τc 1 + ω 2τ c2
[3.3]
Equation [3.2] can be expressed as a function of temperature by assuming a thermally activated process for the lithium motions, where an Arrhenius
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temperature dependence for the correlation time is often assumed (Eq. 3.1). Activation energies are calculated from the linear slope of the Arrhenius plot of the T 1−1 data. Equations [3.2] and [3.3] predict the presence of a spin-lattice relaxation rate T 1−1 maximum at a given temperature, Tmax, at which the condition ω0τc ≈ 0.62 is fulfilled. Therefore, the position of the T 1−1 maximum indicates the temperature at which the motional correlation time τc is comparable to the reciprocal of the Larmor frequency. This is a reasonable approximation for any relaxation model, so that one can be confident regarding evaluations of τc at the observed T 1−1 maximum. It is possible to compare the relative mobility of the lithium in different samples measured at the same ω0 by comparing their Tmax. For example, the shift in Tmax for the electrolyte formed by HEC/DPEO/LiClO4 in Fig. 3.8 toward a higher temperature reflects a reduced lithium mobility in this electrolyte when compared to those based on amylopectin plasticized with glycerol. It should be noted that the NMR relaxation curves observed in plasticized polymer gel electrolytes always exhibit an asymmetric shape around the relaxation rate maximum. This behavior was interpreted by assuming that there are at least two distinct lithium ion dynamics, which were tentatively associated with the complexation of the lithium ion with both the polymer and the plasticizer.87,88 It is interesting to observe the NMR relaxation study of the electrolyte formed by PEO and chitosan complexed with LiClO4. The relaxation data were analyzed by assuming two lithium dynamics; the first one associated with the motion of the lithium ion dissolved in PEO and the second associated with the ions complexed by the chitosan. The experimental 7Li relaxation data were deconvoluted into two theoretical spin-lattice relaxation processes following Eqs. [3.2] and [3.3] (Fig. 3.9). Table 3.5 summarizes parameters τ0 and Ea obtained from the 7Li NMR spin-lattice relaxation data for PEO-based and natural polymer-based electrolytes. Figure 3.10 shows the temperature dependence of the 1H NMR spinlattice relaxation rates (T 1−1) of two proton conducting polymer gel electrolytes. The first is based on a gelatin network and the second on agar, both containing acetic acid plasticized with glycerol and crosslinked with formaldehyde. In this case the relaxation mechanism was the motional modulation of the 1H dipolar interactions and constant C in Eq. [3.2] is related to the mean square amplitude of the fluctuation of 1H–1H dipole-dipole interaction. In both gel electrolytes the 1H spin-lattice relaxation showed a welldefined rate maximum whose position depends on the acid concentration and reflects the high proton mobility in these proton conducting gel electrolytes. In relation to the proton conduction mechanism in these proton gel electrolytes, Mattos et al.97 postulated that the ionic conductivity results from a
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1/T1 (s–1)
101
100
10–1
2.5
3.0
3.5 4.0 4.5 1000/T (K–1)
5.0
3.9 Deconvolution of the 7Li spin-lattice relaxation rates data of the electrolyte formed by PEO and chitosan complexed with LiClO4 in two BPP relaxation processes. The observed relaxation rates are given by the sum (solid curve) of the PEO/Li contribution (dotted curve) and a contribution due to the chitosan/Li (dashed curve). Reprinted with permission from Donoso et al.91
Table 3.5 NMR parameters obtained from the 7Li NMR experiments in polymer electrolytes. Tmax is the temperature at the spin-relaxation rate maximum, τ0 is a pre-exponential factor in the Arrhenius expression, Eq. [3.1], and Ea is the activation energy for the motion causing the relaxation Polymer electrolyte
Tmax (K)
Ea (eV)
τ0 (s)
PEO8LiClO4 HEC/DPEO/LiClO4 Amylopectin/glycerol/LiClO4 PEO/chitosan/Li
350 364 323 330 270 –
0.20 0.24 0.24 0.22 0.28 0.3–0.6
9 3 4.7 4 2 –
Gelatin/glycerol/LiClO4
× × × × ×
Reference 10−13 10−13 10−12 10−13 10−14
80 13 88 91 90
combination of both the vehicular and the Grotthuss-type mechanism. The single line-narrowing transition observed in the 1H NMR spectrum was related to a Grotthuss-like mechanism while the asymmetric shape of the spin-lattice relaxation curve (Fig. 3.10) evoked distinct ion dynamics, which
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1/T1 (s–1)
1
H - 36 MHz
101
100
2.5
3.0
3.5
4.0
4.5
5.0
1000/T (K–1)
3.10 Temperature dependence of the 1H spin-lattice relaxation rates (1/ T1) in the gel electrolytes measured at the Larmor frequency of 36 MHz: (䊉) gelatin network and (䊊) agar network. The electrolytes are formed by the network and acetic acid, plasticized with glycerol and crosslinked with formaldehyde.
is associated with a different complexation of the ions, suggesting the possibility of a vehicle-type mechanism. According to Jeffrey et al.,95 the motion of the protonic species represents the vehicular mechanism for the proton transport, while the proton hopping from one species to others constitutes the Grotthus mechanism.
3.5.3 Electron paramagnetic resonance The EPR spectroscopy was recently used to investigate the coordination environment of the cation in gel electrolytes.96,97 The analysis was performed in electrolytes based on a gelatin containing acetic acid plasticized with glycerol and crosslinked with formaldehyde, doped with copper perchlorate. The X-band EPR spectrum (Fig. 3.11) indicates the presence of the paramagnetic Cu2+ ions in axially distorted sites. The Cu2+ ion has a 3d9 electronic configuration (spin S = ½) and the nuclear spin for both 63Cu and 65Cu isotopes is I = 3/2. The hyperfine structures result from the dipole–dipole interaction between the magnetic moment of the nucleus and the electronic moment of the Cu2+ paramagnetic ion. A closer examination of the line shape in the low field part of the spectrum (parallel region) of Figure 3.11 reveals that the experimental spectrum is a superposition of two spectral compo-
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Intensity (a.u.)
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2500
2750
3000
3250
3500
3750
Magnetic field (gauss)
3.11 X-band CW-EPR spectrum measured at 30 K of the gel electrolyte formed by a gelatin network, acetic acid, plasticized with glycerol crosslinked with formaldehyde and doped with copper perchlorate.
nents, corresponding to two different copper species. The experimental Cu2+ EPR spectrum was analyzed by numerical simulation of an axial spin Hamiltonian, which includes the hyperfine interaction. In order to verify the involvement of the nitrogen atom in the copper complexation of the gel electrolyte, the authors used the pulsed EPR technique electron spin echo envelope modulation (ESEEM). The decay of the electron spin echo due to the interactions between the paramagnetic species and neighboring nuclear spins is usually modulated with frequencies that correspond to the nuclear Larmor ones. The ESEEM results showed the interaction of 1H and 14 N nuclei with the Cu2+ ions, suggesting that both the 1H nuclei (of the gelatin and the plasticizer) and the 14N nuclei (of the glycine and proline groups of the gelatin) interact with the copper ion complexed in the gel electrolyte.96,97
3.6
Conclusions and future trends
Ending this review on electrolytes based on natural polymers we can cite again Michel Armand, who 15 years ago said ‘The research on polymer electrolytes is certainly not on a slowing course, with yet unsettled scientific problems like dissociation and transport numbers which still call for novel ideas, techniques and, surely, patient work’.5 This remark is still valid, even with much more knowledge. As demonstrated in this chapter, among different polymer electrolytes proposed during the last 30 years, there are also SPEs or gel electrolytes
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based on natural polymers. As considered by many researchers, the main obstacle to use natural polymer-based electrolytes in practical applications, as electrochemical devices, is the presence of water, and for this purpose many proposed systems are principally based on PEO and organic, expensive, toxic and less volatile solvents. As shown in this review, polymer electrolyte systems with low price good ionic conductivity values and of easy preparation are those containing plasticizers. However, the major part of these materials is obtained using water as solvent. It was also stated that the presence of water improves the ionic conductivity values, so the question is: is there future for this type of material? From the materials science point of view the answer is no, but if we consider live science and observe our environment we can propose new ideas. Our body is composed of 70% of water and our water-based electrolytic system is very important for our life; most of the environmental systems are biodegradable and necessary to life cycle; the efficiency of some environmental systems, like photosynthesis is very low and the durability depends on the system. Concerning real materials science research, if we consider practical applications, we can cite two of few works published in prestigious journals. Hayes and Feenstra101 presented the video-speed electronic paper based on electrowetting, where one of the important coatings is water. The authors showed that a DC voltage can precisely control the microfluidic motion of colored oil films sandwiched between a solid hydrophobic insulator and water. Moreover this reflective display showed a reflectivity of over 35%, a contrast of 15 and analogue grey scales. Nogi and Yano102 obtained flexible transparent plastic substrates for electrochemical devices with an ultra-low coefficient of thermal expansion (CTE) of 4 ppm/K based on bacterial cellulose nanocomposites (bacterial cellulose contains water equivalent to at least 200 times the dry weight of the cellulose). They also applied this material to organic light-emitting diode and stated that owing to the bacterial cellulose pellicles continuous cultivation process, this material has the potential for electronics display industry, which can use less expensive and continuous roll-to-roll manufacturing process. Considering all these arguments, to use and to give more attention to the natural polymers as potential candidates for the development of new materials we can conclude that ‘The choice of new materials with finely tuned properties appears still almost limitless’.5
3.7
References
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3 fenton, d.e., parker, d.m. and wright, p.v. 1973, Polymer, 14, 589. 4 armand, m.b. 1973, Fast Ion Transport in Solids, W. Van Gool (Eds.), NorthHolland, Amsterdam, 665. 5 armand, m. 1994, Solid State Ionics, 69, 309. 6 gray, f.m. 1991, Solid Polymer Electrolytes, Fundamentals and Technological Applications, VCH, Publishers Inc, New York. 7 munro, b., conrad, p., kramer, s., schmidt, h. and zapp, p. 1998, Sol. Ener. Mater. Sol. Cells, 54, 131. 8 dahmouche, k., atik, m., mello, n.c., bonagamba, t.j., panepucci, h., aegerter, a., and judeinstein, p. 1997, J. Sol-Gel Sci. Techn., 8, 711. 9 dahmouche, k., atik, m., mello, n.c., bonagamba, t., panepucci, h., judeinstein, p. and aegerter, m.a. 1998, Sol. Ener. Mat. Sol. Cells, 54, 1. 10 zoppi, r.a., fonseca, c.m.n.p., depaoli, m.a. and nunes, s.p. 1996, Solid State Ionics, 91, 123. 11 zoppi, r.a. and nunes, s.p. 1997, Polímeros: Ciência e Tecnologia, 4, 27. 12 regiani, a.m., curvelo, a.a.s., gandini, a. and pawlicka, a. 2000, Mol. Cryst. Liq. Cryst., 353, 181. 13 tambelli, c.e., donoso, j.p., regiani, a.m., pawlicka, a., gandini, a. and lenest, j.f. 2001, Electrochim. Acta, 46, 1665. 14 dragunski, d.c. and pawlicka, a. 2002, Mol. Cryst. Liq. Cryst., 374, 561. 15 osman, z., ibrahim, z.a. and arof, a.k. 2001, Carbohyd. Polym., 44, 167. 16 vieira, d, avellaneda, c. and pawlicka, a. 2007, Electrochim. Acta, 53, 1404. 17 glasse, m.d., idris, r., latham, r.j., linford, r.g. and schlindwein, w.s. 2002, Solid State Ionics, 147, 289. 18 lorcks, j. 1998, Polym. Degrad. Stab., 59, 245. 19 singh, j., kaurb, l. and mccarthy, o.j. 2007, Food Hydrocolloids, 21, 1. 20 galliard, t. 1987, Starch: Properties and Potentials, John Wiley & Sons, New York. 21 van soest, j.j.g. and vliegenthart, f.g.j. 1997, Trends in Biotechnology, 15, 208. 22 dragunski, d.c. and pawlicka, a. 2001, Mater. Res., 4, 77. 23 chung, s.h., heitjans, p., winter, r., bzauchac, w., florjanczyk, z. and onodab, y. 1998, Solid State Ionics, 112, 153–159. 24 schoenenberger, c., le nest, j.f. and gandini, a. 1995, Electrochim. Acta, 40, 2281. 25 le nest, j.-f., gandini, a. and schoenenberger, c. 1994, Trends Polym. Sci., 2, 432. 26 regiani, a.m., pawlicka, a., curvelo, a.a.s., gandini, a. and le nest, j.-f. 1999, Polímeros: Ciência e Tecnologia, 3, 45. 27 regiani, a.m., tambelli, c.e., pawlicka, a., curvelo, a.a.s., gandini, a., lenest, j.-f. and donoso, j.p. 2000, Polym. International, 49, 960. 28 regiani, a.m., machado, g.o., le nest, j.-f., gandini, a. and pawlicka, a. 2001, Macromol. Symp., 175, 45. 29 chelmecki, m., meyer, w.h. and wegner, g. 2007, J. App. Polym. Sci., 105, 25. 30 machado, g.o., yonashiro, m. and pawlicka, a. 2004, Nonlin. Optics Quant. Optics, 32, 141. 31 rees, c.w. and storr, r.c. 1968, Chem. Commun., 21, 1305. 32 pillard, d.a., cornell, j.s., dufresne, d.l. and hernandez, m.t., 2001, Water Res., 35, 557. 33 rinaudo, m. 2006, Prog. Polym. Sci., 31, 603.
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65 vieira, d., avellaneda, c.o. and pawlicka, a. 2008, Mol. Cryst. Liq. Cryst., 485, 843. 66 vieira, d. and pawlicka, a. 2010, Optimization of performances of gelatin/LiBF4 based polymer electrolytes by plasticizing effects, Electrochim. Acta, 55, 1489. 67 avellaneda, c.o., vieira, d., alkahlout, a., heusing, s., leite, e., pawlicka, a. and aegerter, m. 2008, Sol. Ener. Mater. Sol. Cells, 92, 228. 68 idris, n.h., majid, s.r., khiar, a.s.a., hassan, m.f. and arof, a.k. 2005, Ionics, 11, 375. 69 majid, s.r. and arof, a.k. 2005, Physica B, 355, 78. 70 khiar, a.s.a., puteh, p. and arof, a.k. 2006, Physica B, 373, 23. 71 fuentes, s., retuert p.j. and gonzález, g. 2007, Electrochim. Acta, 53, 1417. 72 wan, y., creber, k.a.m., peppley, b. and bui, v.t. 2003, Macromol. Chem. Phys., 204, 850. 73 wan, y., creber, k.a.m., peppley, b. and bui, v.t. 2004, J. Appl. Polym. Sci., 94, 2309. 74 sato, t., banno, k., maruo, t. and nozu, n. 2005, J. Power Sources, 152, 264. 75 kim, s.j., shin, s.r., spinks, g.m., kim, i.y. and kim, s.i. 2005, J. App. Polym. Sci., 96, 867. 76 tiwari, a., sen, v., dhakate, s.r., mishra, a.p. and singh, v. 2008, Polym. Adv. Technol., 19, 909. 77 böhmer, r., jeffrey, k.r. and vogel, m. 2007, Progress in NMR Spectroscopy, 50, 87. 78 chadwick, a.v. and worboys, m.r. 1987, Polymer Electrolytes Reviews I, J.R. MacCallum and C.A. Vincent (Eds.), Elsevier Applied Science, New York, 275. 79 chung, s.h., jeffrey, k.r. and stevens, j.r. 1991, J. Chem. Phys., 94, 1803. 80 donoso, j.p., bonagamba, t.j., panepucci, h.c., oliveira, l.n., gorecki, w., berthier, c. and armand, m. 1993, J. Chem. Phys., 98, 10026. 81 forsyth, m., garcia, m., macfarlane, d.r., meakin, p., ng, s. and smith, m.e. 1996, Solid State Ionics, 85, 209. 82 johansson, a. and tegenfeldt, j. 1996, J. Chem. Phys., 104, 5317. 83 mustarelli, p., capiglia, c., quartarone, e., tomasi, c. and ferloni, p. 1999, Phys. Rev. B, 60, 7228. 84 volel, m., armand, m. and gorecki, w. 2004, Macromolecules, 37, 8373. 85 abbrent, s., chung, s.h., greenbaum, s.g., muthu, j. and giannelis, e.p. 2003, Electrochim. Acta, 48, 2113. 86 byrne, n., efthimiadis, j., macfarlane, d.r. and forsyth, m. 2004, J. Mater. Chem., 14, 127. 87 lopes, l.v.s., machado, g.o., pawlicka, a. and donoso, j.p. 2005, Electrochim. Acta, 50, 3978. 88 lopes, l.v.s., dragunski, d.c., pawlicka, a. and donoso, j.p. 2003, Electrochim. Acta, 48, 2021. 89 mattos, r.i., tambelli, c.e., donoso, j.p., and pawlicka, a. 2007, Electrochim. Acta, 53, 1461. 90 mattos, r.i., pawlicka, a., tambelli, c.e. and donoso, j.p. 2008, Mol. Cryst. Liq. Cryst., 483, 120. 91 donoso, j.p., lopes, l.v.s., pawlicka, a., fuentes, s., retuert, p.j. and gonzález, g. 2007, Electrochim. Acta, 53, 1455. 92 przytuski, j., poltarzewski, z. and wieczorek, w. 1997, Polymer, 39, 4343.
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93 wieczorek, w., zukowska, g., borkowska, r., chung, s.h. and greenbaum, s. 2001, Electrochim. Acta, 46, 1427. 94 chung, s.h., wang, y., greenbaum, s.g., bzducha, w., zukowska, g. and wieczorek, w. 2001, Electrochim. Acta, 46, 1651. 95 jeffrey, k.r., zukowska, g.z. and stevens, j.r. 2003, J. Chem. Phys., 119, 2422. 96 mattos, r.i., pawlicka, a., lima, j.f., tambelli, c.e., magon, c.j. and donoso, j.p. 2008, Proc. 11th International Symposium on Polymer Electrolytes, 11th ISPE, Ofir, Portugal, Abstract P-19, page 106. 97 mattos, r.i., pawlicka, a., lima, j.f., tambelli, c.e., magon, c.j. and donoso, j.p. 2010, Electrochim. Acta, 55, 1396. 98 raphael, w., avellaneda, c.o., rodrigues, b.v.m. and pawlicka, a. 2008, Proc. 11th International Symposium on Polymer Electrolytes, 11th ISPE, Ofir, Portugal, Abstract P-58, page 145. 99 klinklai, w., kawahara, s., mizumo, t., yoshizawa, m., isono, y. and ohno, h. 2004, Solid State Ionics, 168, 131. 100 samir, m.a.s.a., alloin, f., gorecki, w., sanchez, j.-y. and dufresne, a. 2004, J. Phys. Chem. B, 108, 10845. 101 hayes, r.a. and feenstra, b.j. 2003, Nature, 425, 383. 102 nogi, m. and yano, h. 2008, Adv. Mater., 20, 1849.
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4 Composite polymer electrolytes for electrochemical devices F. A L L O I N and C. I O J O I U, LEPMI CNRS, Grenoble Institute of Technology, France
Abstract: This chapter discusses composite polymer electrolytes for lithium batteries and proton exchange membrane fuel cell (PEMFC) applications, and describes how nanofillers affect polymer matrix morphologies, dynamic and electrochemical electrolyte performances. The chapter first reviews composite polymer electrolytes for lithium batteries based on organic and inorganic charges, and then discusses composite polymer electrolytes for PEMFC that use ionomer/polymer and hydroscopic inorganic particles to keep the membranes hydrated at higher operating temperatures, and finally self-humidifying membranes, in which water is produced in the membrane. Key words: composite polymer electrolytes, fillers, lithium batteries, PEMFC, electrochemical property, HT-PEMFC.
4.1
Introduction
Composites, i.e. filler and matrix, provide an interesting potential solution to the need for new materials. In composites, each component retains its own properties. In multi-phase materials, the overall properties depend on those of the filler and the matrix, but are also associated with specific interactions between the two materials. In the field of polymer electrolytes for electrochemical applications, the functional properties required are numerous and often incompatible, for example a polymer matrix with a high dynamic that also exhibits good mechanical properties. In the 1980s, several authors proposed the use of composite polymer electrolytes. The solutions they proposed depended on the electrochemical application, i.e. lithium batteries, fuel cells, etc., which determined the properties required. This chapter reviews the development and properties of composite polymer electrolytes used in lithium batteries and proton exchange membrane fuel cells (PEMFC). The effects of fillers on electrolyte properties are discussed in terms of electrochemical performance, and also in terms of polymer matrix morphology and dynamics. Data from the literature are compared in order to determine the effects of the manufacturing 129 © Woodhead Publishing Limited, 2010
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process, particle characteristics (size, nature, amount) and polymer matrix on electrochemical performances.
4.2
Composite electrolytes for lithium batteries: introduction
Polymer-based ion conducting materials have been of great interest to researchers in the field of lithium batteries since Armand et al.1 proposed the use of poly(ethylene oxide) (PEO)–Li salts as a solid polymer electrolyte (SPE). In this application, the polymer electrolyte functions as a mechanical separator between the two electrodes and also as the ionic conductor. Polymer electrolytes are used in the form of thin films and may be either dry (organic solvent-free) or plasticised. A high specific energy density can be reached at medium temperature using a dry polymer electrolyte and lithium metal as the negative electrode. Many SPEs are based on PEO doped with lithium metal salts. PEO-based electrolytes exhibit their highest conductivity values at high temperatures due to their ability to solvate lithium cations. However, the high degree of crystallinity of PEO strongly restricts ionic conductivity at room temperature (melting temperature (Tm) is around 65 °C). Since local relaxations and segmental motions of the polymer host are needed2 for efficient Li+ transport in the electrolyte, ionic mobility occurs predominantly in the amorphous phase of the polymer. Thus it is only above the melting temperature that appreciable conductivities (σ > 10−4 S/cm) are obtained for crystalline PEO–LiX complexes. The polymer melting point may be lowered using a ‘plasticiser’ salt, such as LiTFSI, lithium bis(trifluoromethane sulphonyl) imide. LiTFSI presents a high salt dissociation level due to the strong electron-withdrawing SO2CF3 groups at both sides of the imide anion, even in a low dielectric medium such as polyether. Thus PEO–LiTFSI complexes exhibit the highest ionic conductivity values for SPEs. However, anionic transport occurs more easily due to the weak interaction between the anion and the polymer matrix. The lithium transference number usually falls within the range of 0.1–0.3 in polyether electrolytes. High anion mobility induces a concentration gradient through the electrolyte, which results in increased internal resistance and voltage losses. Research is needed into new salts exhibiting higher transference numbers, such as lithium aluminates bearing PEO chains.3 Many approaches have been used to improve the conductivity of polymer electrolytes. Studies have focused on the development of copolymers4–11 and cross-linked polymers12–14 to create completely amorphous electrolytes with enhanced conductivity at room temperature. One method involves the addition of various plasticisers, such as ethylene carbonate (EC), propylene carbonate (PC) and tetraethylene glycol (TEG), which have been added to
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polymers to form gel polymer electrolytes. These electrolytes have conductivities higher than 10−4 S/cm at room temperature. Unfortunately, most of them are not suitable for lithium metal battery applications as the plasticisers can volatilise, leak, decompose or react with the lithium metal electrode. Furthermore, the addition of liquid plasticisers to polymer electrolytes can destroy the mechanical stability of these electrolytes, which is one of their main advantages. Another strategy for enhancing conductivity is to add inorganic fillers. Since the first study by Weston and Steele,15 which reported improvements in the mechanical properties of (PEO)8LiClO4 by doping with alumina powder (particle size 40 µm), many other studies have been conducted on composite polymer electrolytes. Such composite electrolytes are commonly obtained by dispersing an inert fine ceramic powder in a polymer electrolyte. But despite the considerable interest in composite electrolytes, these materials and the dominant conduction mechanism are still not fully understood. Thus, conflicting data are found in the literature, and data can be difficult to compare since different processes, e.g. drying process, composite electrolyte elaboration, and measurement procedures are used to study the electrolytes. The aim of the first part of this chapter is to present recent research on composite polymer electrolytes, both solid and gel, in order to try to clarify the effect of fillers on ionic conduction, transference number, the polymer crystallisation process, mechanical properties and interface properties in polymer electrolytes.
4.3
Solid polymer electrolytes
Most studies on composite polymer electrolytes have been performed on a PEO matrix. However, PEO is very sensitive to thermal,16 photochemical17 and ultrasound18 degradation. In ultrasound degradation, chain scission is caused by intense shear stresses arising from the collapse of transient cavitation bubbles. The resulting breakage of covalent bonds occurs preferentially in the middle of polymer chains,19 and down to a molecular weight limit of about 2 × 104 g/mol, below which the PEO will not undergo scission. Sonication also produces heat, leading to local increases in temperature and associated thermal degradation, even if the temperature of the sample is controlled externally. Consequently, nanocomposite electrolytes based on PEO require controlled processes. Some authors have used ultrasonic processes20 in order to ensure homogeneous dispersion of fillers in the PEO matrix, but this may induce PEO degradation. The molecular weight of the PEO must be controlled in order to be sure that no degradation occurs, and the polymer matrix must also have the same molecular weight in filled and unfilled electrolytes, so that the real effects of nanoparticle addition can be studied.
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4.4
Composite polymer electrolytes based on poly(ethylene oxide) and clays
4.4.1 Cationic clay Composite polymer electrolyte elaboration One strategy for improving the electrical and mechanical properties of polymer electrolytes is the development of polymer silicate nanocomposites. Following work reported by Usuki et al.,19 a number of studies have been carried out on modifying polymers by incorporating nanostructured layered silicates. The silicates used share the same structural characteristics as the better known talc and mica, and their crystal lattice consists of negatively charged aluminosilicate layers. In their pristine form, this negative charge is balanced by hydrated cations (Li+, Na+ or Ca2+) that occupy the spaces between the layers. PEO can be intercalated20–23 between these layers by replacing some or all of the water molecules occupying the interlayer. With clay and polymer, several types of composite are possible. A classic composite polymer does not enter the clay interlayers. If the polymer diffuses into the interlayers and the layers remain stacked, an intercalated composite is formed. If interplanar attraction is overcome and the silicate layers are separated, the composite is exfoliated. These different morphologies have a large impact on functional properties. Owing to the presence of water in the clay structure, the clay must be dried carefully. Montmorillonite (MMT) can be dried by annealing at 250 °C for about 3 hours to remove water molecules from the interlayers.24 PEO intercalation can be obtained by melt intercalation, in which a dry PEO/silicate mixture is heated above the melting temperature of the PEO.24–26 Alternatively, the nanocomposites can be prepared from solution,27,28 resulting in nanocomposite structures similar to those prepared by the melt intercalation process. In order to accelerate PEO intercalation, Aranda et al.29 formed the nanocomposite in melt state with the use of microwave irradiation. The process must be controlled carefully to avoid PEO degradation due to localized high temperatures. The incorporation of PEO leads to clay agglomeration with a disorganised stacking arrangement. Structure modification Most results in the literature concern nanocomposites based on PEO and Li+-exchanged clays, i.e. the cations originally present in the intergallery spacing, usually Na+, have been exchanged for Li+.28–32 Kurian et al.27 studied the formation of polymer–silicate nanocomposite electrolytes based on lithium cation exchanged montmorillonite clay, Li-MMT, or organically
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modified Li–MMT and an amorphous polyether, poly(oxyethylene)8 methacrylate, POEM. The intercalation of POEM in Li–MMT was characterised by X-ray scattering. An increase in the spacing (001) planes from 10 Å in the native Li–MMT to between 17 and 18 Å was observed, depending on the amount of clay introduced in the polymer. At lower clay loadings, around 5 wt%, in addition to intercalated regions, the presence of individual delaminated clay indicated that some exfoliation had occurred in the nanocomposites. At higher loadings, around 15 wt%, a limited amount of exfoliated clay and a possible phase separation were observed. The nonuniform dispersion of the clay in the polymer matrix may form physical barriers detrimental to ion motion. In an attempt to promote more uniform dispersion of the clay at higher loadings, nanocomposites incorporating organically modified Li–MMT were evaluated. The organic Li–MMT clays were obtained by partial modification of the Li–MMT cation exchange site with (NH2)2PEO.26 The incorporation of organic Li–MMT clays in the PEO improved the dispersion of the clay. Ratna et al.34 characterised PEO/clay nanocomposite electrolytes using an apolar Cloisite® nanoclay modified with alkylammonium cations. Modification of the clay with organic ions not only makes the clay surface hydrophobic, but also results in a very large increase in d-spacing, which may facilitate the incorporation of PEO into the interlayer galleries. Intercalated nanocomposites were obtained.34 Shen et al.24 studied nanocomposites obtained by melt intercalation based on PEO/Na+–MMT and PEO/B34 (bentonite with dimethyl ditallow ammonium cations).24 The presence of water molecules in the interlayers of MMT did not have any significant effect on the melt intercalation process, i.e. saturation ratio and final layer distance in the composites. In a hybrid composite PEO/B34/Na+–MMT, the PEO first intercalated into B34, and later into MMT. Polymer molecular weights did not affect the nanocomposite structure or saturation ratios of polymer to layered silicates, but did affect the kinetics of melt intercalation.24 The influence of both PEO molecular weight and the nature of Cloisite®, Cloisite®-Na+, Cloisite® 20A (organically modified with dimethyl dihydrogenated tallow quaternary ammonium) and Cloisite® 30B (organically modified with methyl tallow bis-2-hydroxyethyl quaternary ammonium cations) on the intercalation process using the melt technique were evaluated.35 Nanocomposites based on Cloisite® 30B, containing a polar modifier, displayed exfoliated nanostructures; the compatibility of the modifier with PEO aided the formation of nanostructure composites as previously noted.26 In contrast, the nanostructure of the Cloisite® 20A-based nanocomposites, containing an apolar modifier, was found to be very dependent on PEO molecular weight. A high molecular weight PEO, and thus a high melt viscosity, was needed to obtain exfoliated structures.
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Chen and Evans36 evaluated the influence of PEO molecular weight on the intercalation process for MMT clay. The clay preferentially absorbed the high molecular weight PEO fractions. This preference may explain the considerable reinforcement observed with MMT clay addition, and raises questions about the effect of the low molecular weight PEO which was excluded. PEO is also able to form intercalates with Laponite®, leading to a basal spacing typical for nanocomposites of PEO and clays with small cations, e.g. Na+ and Li+, in the interlayer spacing.30 But the lack of an organic modifier in the Laponite® interlayer space prevents the formation of highly intercalated nanostructures. X-ray results indicated that the maximum amount of PEO that can be intercalated into Laponite® is slightly less than 0.7 g of PEO/g of Laponite®. Greater polymer amounts induced the formation of two-phase mixtures.30 The presence of salt also has a considerable effect. The addition of Cloisite® Na+ and Laponite® resulted in the formation of intercalated PEO nanocomposites, without NaClO4.37 However, NaClO4 containing PEO/Cloisite® nanocomposites displayed an exfoliated nanostructure. The presence of exfoliated, single silicate layers results in a high effective aspect ratio, which may have a large impact on the properties of nanocomposite polymer electrolytes. Thus, it would appear to be possible to control nanocomposite structure through the choice of clay type, the presence or absence of an organic modifier, and the salt concentration. In these nanocomposite electrolytes, PEO is confined in well-defined slits, formed by negatively charged parallel stacks. In general, the added clay does not act as a nucleating agent for PEO crystallisation, as expected from the decrease in both crystallisation temperatures and crystallinity when compared with pure PEO.35,38,39 The presence of the dispersed clay disturbs PEO crystal growth due to the restricted chain mobility in the limited space available. The decrease in crystallisation rate does not appear to be directly related to the incorporation of the polymer chain in the interlayer structure, as no distinct difference in crystallisation behaviour was observed between intercalated (Cloisite®-Na+ and Laponite®) and exfoliated (Cloisite® 30B) PEO nanocomposites.35,40 The decreases in melting temperature and crystalline phase content of the polymer matrix are associated with the formation of less perfect crystalline structures. In the presence of salt, the effect of clay on PEO crystallisation depends on the nature and amount of clay and on the salt concentration.41 At high salt concentrations, the nanocomposite electrolyte PEO/NaClO4 based on Cloisite® 30B exhibited a higher melting point and crystalline content than unfilled electrolyte. Chen and Chang42 observed the same behaviour, i.e. an increase in PEO crystallisation, for PEO nanocomposites based on MMT modified with poly(oxypropylene) diamine and LiCF3SO3 salt, for a small amount of clay. Increasing the
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amount of clay induced a decrease in PEO crystallisation. The authors explained the increase in crystallisation by the interaction between the cation of the lithium or sodium salt with the clay, which decreased PEO/ cation interaction. This weaker interaction increased PEO flexibility, thus enabling crystallisation. Conductivity behaviour Single-ion conductors can be obtained by the intercalation of PEO on clay due to the presence of cation charge at the silicate surface.27,29 The conductivity values of electrolytes based on POEM with the addition of 2 and 5 wt% clay were found to be around 4 × 10−7 S/cm at 70 °C.27 The conductivity obtained can be anisotropic.30 Molecular dynamic simulation has shown that the Li+ ions are solvated preferentially by the silicate oxygen atom rather than PEO.43 The conductivity is too low for practical applications, even with a cationic transference number equal to one. In order to increase conductivity, but with a cationic transference number different from one, lithium salts were added to PEO/clay nanocomposites. At room temperature, the nanocomposite electrolyte exhibited higher ionic conductivity than unfilled polymer due to the larger content of the PEO amorphous phase.28,42 The improvement in conductivity depends on the nature of the clay. Fan et al.28 have shown that 250-Li–MMT, i.e. Li–MMT heated to 250 °C, was more effective in enhancing the conductivity of (PEO)16LiClO4 than Org–MMT, dodecylamine modified Li–MMT, and Li–MMT, since 250-Li–MMT forms an exfoliated structure in the PEO matrix. The maximum conductivity at room temperature was obtained with a small amount of clay, generally between 3% and 5 wt%.42 Larger amounts of clay caused a blocking effect due to the incorporation of poor conducting rigid particles. At high temperatures, the unfilled electrolyte exhibited the highest conductivity, even if the conductivity difference was small28,37 (Fig. 4.1). This was due to the presence of clay platelets, which hindered ion transport. Using computer simulations, Hackett et al.44 determined that the cations resided near the silicate surface rather than being coordinated with PEO in dried nanocomposite PEO/clay electrolytes. In hydrated nanocomposites, more than half of the cations were near the silicate surface. Ions near the silicate surface move significantly more slowly than those in the centre of the interlayers. Cations are coordinated by water and, to a lesser extent, by the silicate surface. In a hydrated system, PEO does not enter significantly into the coordination cell of the cations. The limitation to ionic diffusivity in these electrolytes may be inefficient coordination of the Li+ with the polyether. Ion dynamics appear to involve sticky hops along the oxide layer of the silicate,45 thus it is very important
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Log (σ) (S/cm)
–4 –6 –8
–10 2.4
PEO/laponite PEO/MMT-30B PEO/MMT-Na+ PEO 2.6
2.8
3
3.2
3.4
3.6
3.8
–1
1000/T (K )
4.1 Conductivity for the PEO/clay (2.5 wt%)/NaClO4 nanocomposites, salt concentration EO/Na+ = 20 from Loyens et al.37
to dry the clay or the electrolyte carefully in order to ensure the conduction behaviour of a solid electrolyte. Mechanical properties The mechanical properties of PEO were improved with the incorporation of clay. Mechanical improvement was associated with the aspect ratio of the clay, with a high aspect ratio being necessary.37 The incorporation of Laponite® in PEO resulted in poor mechanical improvement due to a low aspect ratio, L/d = 25 (L being the length and d the diameter).40 Nanocomposites with an exfoliated structure exhibited the best mechanical properties.35 Kurian et al.27 observed an increase in the modulus of the POEM matrix when the clay content was increased by six orders of magnitude at 20 wt% clay. The loss modulus G″ was greater than G′ at lower clay loading; at higher clay loading the opposite trend was observed, indicating a switch from liquid-like to solid-like behaviour.
4.4.2 Anionic clay Inorganic layered materials with positive charges, such as layered double hydroxides, LDH, can be used in nanocomposite polymer electrolytes.43 Experiments were performed on LDHs composed of positive charged layers MIIxMIII1−x(OH)2 and interlayer exchangeable anions (Am−x/m · nH2O), where A was the anion, and MII and MIII were divalent and trivalent metal cations.46 An intercalation or an exfoliation nanocomposite is obtained by mixing PEO and cationic clay. PEO cannot easily intercalate into the inter-
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layers of LDH. In fact no specific interaction exists between the anions in the interlayers and the PEO. Yang et al.47 reported polyoxyethylene sulphate intercalation in LDH. This PEO modification involves the presence of anions at the PEO end-groups. A specific interaction may then occur allowing PEO to intercalate. The structure obtained depends on the nature of the clay. In PEO composites, an exfoliation morphology was obtained using oligo(oxyethylene)-modified LDH, OLDH,43,48 whereas an intercalation morphology was obtained with pristine LDH.43 The intercalation of PEO with an acid end-group was obtained by Leroux et al.49 An IR study indicated a strong interaction between the PEO-sulphate or the PEO carboxylic acid oligomers and the layer surface. The addition of OLDH was found to increase the PEO melting temperature, with the exfoliated dispersed OLDH acting as a nucleating agent with a strong heterogeneous nucleation effect on PEO crystallisation.48 At 20 °C, the conductivities of nanocomposite electrolytes PEO/LiClO4/OLDH and PEO/LiClO4/LDH were 1.6 × 10−5 and 2.8 × 10−6 S/cm, which were higher than the unfilled sample. This improvement was related to a decrease in the amount of crystallinity. The higher conductivity of PEO/LiClO4/OLDH compared with PEO/LiClO4/LDH can be explained by a good dispersion of compatible LDH layers in the PEO matrix. At high temperatures, the conductivities of the filled and unfilled electrolytes were very similar.48 The TLi+ values measured at 90 °C were 0.21, 0.42 and 0.30 for pure PEO/LiClO4, PEO/LiClO4/LDH and PEO/LiClO4/OLDH nanocomposites respectively.48 Owing to the Lewis acid OH groups located on the layer surface, a strong Lewis acid–base interaction between the LDH surface and the ClO4− anions may occur, and a decrease in anion mobility and an improvement in the cationic transference number may be obtained.
4.5
Composite polymer electrolytes based on poly(ethylene oxide) and non-ionic fillers
4.5.1 Metal oxide charges Composite polymer electrolyte elaboration Since the first studies performed in the 1980s, a large number of metal oxide charges have been added to polyether matrices.50–79 The fillers used are generally SiO2,55–58,76,78 ZnO,73,77,79 TiO2,59,61,67,75 Al2O350,59–67,72,76 and γAlO2.51– 54,69 The main process used to obtain nanocomposite polymer electrolytes is the film casting technique. The nanoparticles and the PEO or the nanoparticles, the PEO and the salt, are mixed together in a common solvent, and a film is obtained by solvent evaporation. Nanocomposite polymer electrolytes can also be obtained by melt mixing and hot pressing,80 and in situ
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nanoparticle sol–gel elaboration can be used to prepare PEO-based nanocomposite polymer electrolytes with uniform dispersion of the fillers.75 Structure modification A net decrease in PEO crystallinity was observed with the incorporation of fillers.59,74,80 In PEO/LiTFSI complexes, an amorphous structure, stable for several months, can be obtained with the incorporation of small amounts of fillers.71 Using 7Li nuclear magnetic resonance (NMR) investigation, Dai et al.67 found that the addition of nanometric Al2O3 to PEO/LiI electrolyte suppressed the formation of crystalline phases. The ceramic fillers, because of their large surface area, may prevent local PEO organisation and significantly reduce PEO crystallisation kinetics. The influence of incorporated nanoparticles on the polyether glass transition temperature,Tg, is open to greater debate. Several studies indicate a quasi invariance of the polyether Tg during the addition of nanofillers. Indeed, no change in Tg was observed for PEO8 LiBF4 with the addition of 10 wt% Al2O3,81 for PEO/LiTFSI with the incorporation of SiO2,71 for PEO8LiClO4 with 20 wt% Al2O382 or for 3PEG, a trifunctional amorphous polyether, with the addition of SiO2 and Al2O3.83 However, a small decrease in PEO Tg was observed with the addition of less than 2 wt% SiO2 to PEO16LiClO4.84 With larger amounts, a Tg increase was observed, although the Tg of the filled PEO remained lower than that of the pure matrix.78 Choi and Shin85 observed a Tg decrease for PEO with the incorporation of SiC and Si3N4. Significant decreases in Tg and viscosity were reported with the incorporation of Al2O3 in polyethylene glycol monomethylether (PEGMME) (Mn = 350 g/mol) for a salt concentration higher than C > 0.5 mol/l.79,86–88 A decrease in Tg was also obtained with the incorporation of acid zirconia, SO42−–ZrO2 in PEO.89 However, the acidic sulphate function is very aggressive for PEO and this reactivity may explain such behaviour. Indeed, Chauvin et al.90 have shown that PEO is very sensitive to an acid medium and a strong decrease in its molecular weight was observed. In other studies, the opposite trend has been observed: a small increase in Tg (one or two degrees) was noted with the incorporation of SiO2 in PEO8LiTFSI and PEO8LiClO4,58 while a large increase in Tg, i.e. 20 °C, and a large decrease in the tan δ peak size obtained by dynamic mechanical analysis (DMA), were observed with the incorporation of hydrophilic SiO2 in PEO/LiClO4.76 The tan δ decrease was associated with a small amount of product involved in the glass transition. A net increase in Tg was observed with the addition of Al2O3 in PEO and in polyoxypropylene.91,92 The Tg of a polymer is directly related to the macroscopic viscosity of the amorphous phase. In a semicrystalline polymer, the amorphous phase is constrained by the crystalline phase, thus the Tg measured in a semicrystal-
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line polymer is higher than the Tg of the same polymer when totally amorphous. Thus, a decrease in the amount of crystalline phase induces a decrease in Tg. It is well known that the incorporation of lithium salts leads to an increase in the Tg of a polyether due to the specific interaction between the lithium cation and the oxygen groups of the polyether. Changes in the Tg of PEO in nanocomposite electrolytes may thus be related to the reduced crystalline phase and the modification or establishment of specific interactions between the polymer, the fillers and the salt. In most of the studies,58,71,81– 84 the effect of filler incorporation on Tg is small, which may be related to antagonist and often weak effects, i.e. decrease in Tg due to PEO amorphisation, specific interactions between fillers and PEO that may increase Tg, or specific interactions between ionic species and fillers that may reduce the cation/polyether interaction and thus decrease PEO Tg. These specific interactions can be investigated using IR and Raman spectroscopy, NMR or quasielastic neutron scattering. Several studies have been carried out on nanocomposite polymer electrolytes using these techniques. Restriction of polymer matrix mobility due to the presence of nanofillers, via filler/polymer interactions, was observed by Tsagaropoulos and Eisenberg93 in nanocomposites with silica aerogel and poly(vinylalcohol), poly(styrene), poly(methylmethacrylate) or poly(4-vinylpyridine) as the polymer matrix. The presence of a second Tg was observed when the silica content was higher than 10 wt%. This Tg, which was higher than the bulk polymer Tg, was associated with the formation of tightly bound polymer, due to polymer/filler interactions. Mobility restriction was also observed by quasi-elastic neutron scattering.94 Elastic scattering was affected by filler addition, with the results obtained suggesting that an immobilised layer of the polyether matrix, 3PEG, was present around the TiO2 filler, roughly 5% of the total polymer with 10 wt% TiO2 (size 21 nm), whereas the bulk polymer was unaffected by the presence of fillers. While 7Li NMR has provided information about ion dynamics, 1H and 13 C NMR have contributed to elucidate PEO dynamics in pure and composite materials. Numerous NMR studies on nanocomposite polymer electrolytes have been published.60,95–100 Owing to the complexity of the phenomena occurring in nanocomposite polymer electrolytes, different observations and interpretations are found in the literature. It is generally expected that the NMR linewidth of the mobile nucleus decreases when ionic motion occurs. However, Dai et al.,60 with highly concentrated PEO1.5/ LiI/Al2O3, and Forsyth et al.,95 with 3PEG/LiClO4/TiO2, SiO2 and Al2O3 nanoparticles, observed an increase in the 7Li linewidth with the addition of nanoparticles, although the ionic conductivity increased. Chung et al.96 showed that the filled electrolyte, PEO8/LiClO4/TiO2, and the unfilled one, exhibit comparable linewidths for nuclei, proton and
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lithium. The 1H diffusion coefficient was slightly decreased with the incorporation of TiO2 at 80 °C. In contrast, enhancement of the 7Li diffusion coefficient was observed with the incorporation of TiO2. Cheung et al.97 also observed an improvement in the lithium diffusion rates without a corresponding increase in segmental motion for PEO/Al2O3/LiCF3SO3. The effect of nanoparticle size was not considered significant, as the highest 7Li diffusion coefficient was obtained with the highest nanoparticle size, 0.05 µm. Mustarelli et al.,98 using PEO8/LiClO4 or LiN(CF3SO2)2/SiO2, found an increase in the activation energy for lithium motion and concluded that 7 Li spin dynamics were influenced by fillers. Using a 7Li–19F decoupling experiment on PEO/LiBF4/α-Al2O3, Bloise et al.99 observed that the Li–F interaction was weaker in the composite with α-Al2O3 than in the composite with γ-Al2O3. The 7Li linewidth decreased significantly when Al2O3 was added to PEG (polyethyleneglycol, Mw = 2000 g/mol)/LiClO4 (O/Li = 26), while there was only a marginal decrease in 1H linewidth.100 Vogel et al.,101 using 2H NMR, found that the presence of TiO2 nanoparticles hardly affected the dynamics or crystallisation behaviour of the polymer in low molecular weight PEO nanocomposites. Bloise et al.102 showed the mobility of Li+ in PEO8LiClO4/TiO2 to be very high. These different behaviours may be associated with the complexity of nanocomposite polymer electrolytes. One possible explanation offered was that nanofillers cause a decrease in Li–H distance and thus an increase in the Li–H heteronuclear dipolar interaction, which is shown to be the main contributor to the linewidth 1H in decoupling experiments.99 Positron annihilation lifetime spectroscopy data suggest that there may be competition between the polymer/Li interaction and the polymer/TiO2 interaction in the composite electrolyte.95 Salt dissociation and PEO conformation depend on the nature and concentration of the salt and the presence of fillers and temperature, and this has been widely studied using IR and Raman spectroscopy.103,104 The D-LAM (disorder-longitudinal acoustic mode) bands between 200 and 300 cm−1 have been assigned to longitudinal backbone motions of the polymer chain.105 The intensity of the PEO D-LAM bands has been shown to decrease significantly with the addition of fillers.106 With the addition of TiO2 to the pure polymer, 3PEG, these modes were suppressed,83 suggesting that there was a restricted mobility of the polymer phase, which agrees with differential scanning calorimetry (DSC) measurements93 and quasi-elastic neutron scattering results.94 In contrast, the addition of Al2O3 did not suppress the D-LAM mode, indicating that the level of interaction between polymer and Al2O3 is not the same as that between polymer and TiO2. The addition of TiO2 did not affect any of the PEO conformation sensitive modes in the CH spectral regions (800–900 cm−1, 2800–3100 cm−1) in 3PEGbased electrolyte.106 In contrast, using a polymer electrolyte based on high
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O C M
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C
M
M
O
M
O
R
O Bidentate
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R
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Scheme 4.1 Three modes of bonding structures from Xiong et al.73
molecular weight PEO and LiClO4, Chen-Yang et al.78 observed a sharpening of the peak associated with the stretching vibrations of the C—O—C when silica aerogel was added. They linked this behaviour to the presence of Lewis acid–base interactions between the silica oxygen atom and the Li+, resulting in a weakening of the interaction between the Li+ and the PEO. A large change in the IR spectrum of PEO/LiClO4 was found to occur with the addition of ZnO.73 The C—O—C stretching vibration becomes a single band centred on 1107 cm−1. The authors linked this modification to the formation of crosslinks that further weaken the C—O bonds. Some evidence of amorphisation of the PEO matrix with the addition of ZnO was also observed using the CH2 wagging mode modification. A change in the coordination modes of acetate groups with zinc metal was observed with its incorporation in the PEO/LiClO4 electrolyte.73 The coordination of lithium ions with both PEO segments and ZnO acetate groups changed the bonding structure of the acetate groups, with bidentate and bridging modes73 (Scheme 4.1). Consequently, the filled electrolyte was more amorphous than the unfilled one. Using the symmetric stretching mode of ClO4 or SO3, no significant decrease in ion pairing was observed with the addition of both TiO2 and Al2O3.83 However, the addition of TiO2 to the electrolyte shifted the symmetric stretching vibrations to smaller frequencies for both triflate and perchlorate salts, indicating weak coordination between the anion and the TiO2 surface. Such behaviour was not observed with Al2O3 fillers.83 Johansson et al.71 studied the interactions between fillers and LiTFSI by directly mixing fillers and salt together. The peaks in the 720–770 cm−1 region revealed no major changes in the anion modes, in particular no increase in the amount of free ions. In order to assess the amount of ion pairing, the ratio between the intensity of the 741 cm−1 band (associated with free ions) and the integrated intensity at 726–759 cm−1 was used.107 An increased ratio was observed with the addition of a small amount of SiO2 to PEO9LiTFSI.71 When the filler amount increased, this ratio decreased, with an associated increase in electrolyte viscosity, which should lower anion mobility and induce broadening owing to the slow orientation of the anion. No direct effect of filler addition on anion dissociation was observed. A small increase in the quantity of free ClO4− ions was obtained with the incorporation of silica aerogel in a highly concentrated PEO6LiClO4 electrolyte.78 The highest
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dissociated electrolyte was obtained with 2% of fillers, subsequently dissociation decreased with an increase in filler amount.78 The dielectric study of nanocomposite polymer electrolytes performed by Jayathilaka et al.108 indicated that the quantity of free ions was not modified significantly with the addition of Al2O3 to PEO/LiTFSI electrolyte. However an increase in ion mobility was noted. In the various studies reported, little or no improvement in ionic dissociation was found, but some specific interactions were observed which induced PEO amorphisation in semicrystalline electrolytes and restricted PEO mobility in amorphous electrolytes. These behaviours were dependent on the nature and amount of filler. Conductivity behaviour Considerable work on improving ionic conduction in nanocomposite polymer electrolytes has been conducted. The best results were obtained with Al2O359,83,86–88,108 and TiO2 ,59,83,96 while improvements with SiO2 were minor.58,109 In PEO electrolytes, conductivity values are very sensitive to the amount of crystalline phase present, which is related to the thermal history of the sample. A large difference in PEO/LiClO4 conductivity at 25 °C was obtained in relation to the sample’s thermal history.73 High conductivity was obtained just after sample cooling, with σ = 10−6 S/cm, then conductivity fell dramatically after the electrolyte had been kept for two weeks at room temperature in a vacuum, σ = 10−9 S/cm. It is difficult to compare conductivity data obtained at room temperature; however, conductivity is generally improved below the PEO melting point after the sample has been previously heated. This improvement is associated with a decrease in the amount of electrolyte in the crystalline phase present59 and a reduction in crystallisation kinetics due to the incorporation of fillers. At room temperature, conductivity was found to increase with the addition of fillers, generally reaching a maximum between 5 and 15 wt% of filler.74,75,97,110–112 For large amounts of filler, conductivity decreased, which was associated with the blocking of conduction pathways by the electrically insulating nanoparticles. The results for conductivity values at high temperature are more conflicting. Some improvement in ionic conductivity has been observed,72,87,97,108,111,112 but many studies show a quasi-invariance in conductivity values at high temperatures,73,106,113 while others report a conductivity decrease.70,71 Croce et al.72 and Jayathilaka et al.108 observed that the nanoparticle surface groups influence improvements in conductivity. The best conductivities were obtained using neutral and acidic Al2O3 nanoparticles (Fig. 4.2). The model developed suggested that enhanced
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–3.0
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Ceramic-free Acid Neutral Basic
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1000/T (K–1)
4.2 Conductivity of PEO20LiCF3SO3 10 wt% Al2O3 nanocomposite electrolytes with different Al2O3 surface groups, extracted from Croce et al.72
Acidic HO Al HO Al
Li+ CF3-SO3–
Al OH Al OH
O Li+ O
CH2 CH2
Neutral Li+
CF3-SO3–
HO Al O Al
Al Al Al
O O
Li+ O
CH2 CH2
Basic Li+
CF3-SO3– O
Al Al
Al Al
Al
Al
O O
Li+
CH2 CH2
O
Scheme 4.2 Model of the surface interactions between nanosized Al2O3 particles and the PEO–LiCF3SO2 electrolyte, extracted from Croce et al.72
ionic conduction is related to the presence of hydrogen bonds between the acidic or neutral ceramic surface and the anion and polymer chain (Scheme 4.2), which improve ionic dissociation. With basic surface nanoparticles, no improvement in ionic conductivity was obtained. A model based on acid– base interactions has been proposed by several authors62,64,72,108 to explain conductivity improvement. However, no evidence of an increase in salt dis-
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Log (σ) (S/cm)
–2.6 5.8 nm
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10–20 nm 37 nm
–3.0 –3.2 –3.4 –3.6 –50
<10 μm
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50 100 150 Surface area of Al2O3 (m2/g)
200
4.3 Evolution of ionic conductivity at 90 °C with specific area of alumina grains for PEO9LiCF3SO3 + Al2O3 composite electrolytes.74
sociation was observed by spectroscopic investigations into similar samples.83 Such interactions may also create ion conducting pathways at the filler surface.59,95 Best et al.83 developed another model based on electrostatic phenomena. Owing to the difference in dielectric constant between the PEO and the fillers, polarisation may occur around the fillers. Such a strong interaction may restrict polymer mobility, which is in accordance with the results of quasi-elastic neutron scattering experiments.94 The cation, due to its localised charge, will show a relatively stable potential landscape, meaning that the potential at the ceramic surface will be of the same order as that at the polymer. The cation will move as the potential barrier is lowered by the filler. Improved ionic conductivity, in the entire temperature range, was observed with the addition of nanoporous Al2O3 (particle size 104 µm, pore size 5.8 nm) to PEO9LiTFSI electrolyte.108 The best conductivity was obtained with acidic Al2O3, with a conductivity of 2 × 10−3 S/cm at 90 °C compared with 6 × 10−4 S/cm for the unfilled electrolyte. Johansson et al.71 showed the opposite behaviour, with a slight decrease in ionic conductivity for both PEO10LiTFSI and PEO40LiCF3SO3 with the addition of 5%wt γ-Al2O3. However, the conductivity of the unfilled sample reported by Jayathilaka et al.108 was lower than that reported by Johansson et al.71 This difference may explain the opposite trends observed. Dissanayake et al.74 observed an increase in conductivity for PEO9 LiCF3SO3 with the addition of Al2O3. The enhancement increases with increasing specific surface area of filler (Fig. 4.3). This evolution is consistent with the interpretation that Lewis acid–base type surface interactions are largely responsible for the observed increase in conductivity.
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–3
Log (σ) (S/cm)
–4 –5 –6 –7 –8 –9 –10 2.7
PEO8–LiClO4–ZnO PEO8–LiClO4 2.8
2.9
3.0 3.1 1000/T (K–1)
3.2
3.3
3.4
4.4 Conductivity versus temperature of PEO8–LiClO4 and PEO8–LiClO4– ZnO (5 wt%).73
An effect of nanoparticle size on ionic conductivity was also observed with the addition of Al2O3 to PEO (2 × 105 g/mol)/LiCF3SO3.97 The best conductivity was observed at high temperature and with 0.05 µm size particles. A small improvement in ionic conductivity was also obtained, at high temperature, by the addition of mesoporous silica to PEO12LiClO4 electrolyte.111,112 A small improvement in ionic conductivity was observed for 3PEG + 1.5 mol/kg LiClO4 at high temperature,106 while no improvement was noted at a lower salt concentration. The same behaviour was reported by Wieczorek et al.,64 with an improvement in the conductivity of a liquid electrolyte based on PEGMME and LiClO4 with the addition of α-Al2O3. This enhancement at high temperatures was only observed for concentrated electrolytes, Csalt > 2 mol/l. At low salt concentrations, the filled and unfilled electrolytes exhibited the same conductivity behaviour at high temperatures. This behaviour was consistent with the idea that the filler influences ion aggregation in cases where salt aggregation is significant. In contrast, a decrease in ionic conductivity over the entire temperature range was observed with the incorporation of SiO2 in a liquid electrolyte based on PEG dimethylether (Mw = 500 g/mol). This decrease was associated with an increase in the viscosity of the electrolyte. The nanocomposite electrolyte PEO–LiClO4–ZnO exhibits the same conductivity as the unfilled electrolyte at high temperatures,73 whereas at room temperature, nanocomposite electrolytes have a much higher conductivity (Fig. 4.4). The nature and amount of fillers were found to have no effect on conductivity at high temperature in PEO–LiSO3CF3 electrolytes obtained by hot-pressing with the addition of both SiO2 and Al2O3 fillers.113
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Cationic transference number and interface properties Even if there is no conductivity improvement at high temperature, the incorporation of fillers can be useful if it improves the cationic transference number or interfacial stability. An increase in cationic transference number, T +, is generally observed with the addition of nanofillers.59,72,77,80,111,114,115 Cationic transference numbers were determined using the electrochemical method developed by Bruce et al.116 This method assumes an ideal electrolyte and complete dissociation of the electrolyte, which is not the case in concentrated electrolytes. However, no method adapted to concentrated electrolytes exists, thus the data must be used with caution. An increase in T + was obtained by the addition of Al2O3 to P(EO/ MEEGE)–LiAl(HFIP)4 (MEEGE = 2-(2-methoxyethoxy)ethyl glycidyl ether, LiAl(HFIP)4 = lithium tetra(1,1,1,3,3,3-hexafluoro-2-propyl) aluminate).114 The highest T + improvement was obtained with the largest amount of Al2O3, even though T + remained low, i.e. 0.12 for 20 wt% Al2O3. An increase in T + was obtained in PEO–LiClO4 electrolytes with the addition of TiO2,59 and mesoporous SiO2.111 In PEO–LiCF3SO3, an improvement was obtained with ZrO2,115 and with acidic and neutral Al2O3 but not with basic Al2O3.72 A moderate decrease in interfacial impedance was observed with the addition of fillers by Appetecchi et al.117 Careful attention must be paid to the elaboration process in order to evaluate the real effect of nanofillers. Xie et al.92 observed no change in interfacial resistance with the addition of SiO2 in high molecular weight PEO, whereas a decrease was observed when a viscous polyether was used. The nanocomposite PEO/LiCF3SO3/γ-LiAlO2 presented good interfacial stability versus metallic lithium.80 However, the TiO2 particles were found to present poor stability versus metallic lithium,50 so that preference must be given to other nanoparticles. Indeed, elaboration conditions, and especially water content, have a major impact on how filler addition affects interfacial properties. With lithium electrodes, cycle duration was much longer with the addition of 20 wt% γ-LiAlO2 to PEO20LiBF4.80 The fillers did not have a direct effect on dendrite growth but the improved mechanical properties of the electrolyte may explain this longer cycle duration.92 Good cyclability in a Li/ PEO35LiCF3SO3 + 5 wt% SiO2/LiFePO4 battery was obtained113 with a capacity fade of only 0.17% per cycle at 100 °C.
Mechanical properties Some authors92,114,118 observed improved mechanical properties in polymer electrolytes with added nanofillers. The rheological behaviour of filled and unfilled PEO–LiTFSI electrolytes was studied at 85 °C.92 A significant
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increase in modulus was obtained with the incorporation of 10 wt% of fumed silica, which may be due to hydrogen bonding between the PEO matrix and the nanoparticle surface. The filled matrix exhibited longer relaxation times, associated with a slower segmental motion. The mechanical properties were evaluated by the dimensional stability of the electrolyte in relation to creep. The latter was indirectly determined by impedance spectroscopy. Improved mechanical stability was reported for α-Al2O3,15 alumina whisker reinforced PEO8–LiClO4118 and hyper-branched polymer electrolyte filled with α and γ-LiAlO2 particles.119 An improvement in tensile modulus above the matrix melting point was observed with the addition of 10 wt% Al2O3.114 Above 50 °C, the unfilled electrolyte P(EO/MEEGE)– LiAl(HFIP)4 flowed, whereas the filled electrolyte exhibited a storage modulus higher than 10 MPa. The storage modulus was stabilised by the incorporation of 10 wt% TiO2 in PEO30LiTFSI, with a storage modulus equal to 4 MPa.120 Croce et al.59 reported a significant increase in both Young’s modulus and yield stress for Al2O3 and TiO2 reinforced PEO8LiClO4. Unfortunately, no modulus value was reported in their paper.
4.5.2 Organic fillers Cellulose whiskers were added to PEO–LiTFSI electrolytes to improve the mechanical properties of the polymer electrolytes without compromising ionic conduction.120–124 The use of cellulose whiskers in nanocomposite polymers was shown to result in a high reinforcing effect at T > Tg and an increase in thermal stability, even at low loading levels. This effect was attributed to the formation of a percolating network within the matrix, resulting from strong hydrogen bonds between nanoparticles. The average aspect ratio, L/d, of these whiskers was estimated to be close to 70. The incorporation of cellulose whiskers in PEO was found to have no effect on PEO Tg, and a decrease in PEO melting temperature was observed only when the whiskers exceeded 10 wt%.122 The same behaviour was obtained with PEO–LiTFSI electrolytes.120 At high temperatures, the storage modulus of the filled electrolyte was found to be greater than that of the unfilled polymer electrolyte by more than a factor of 100.120 Tunicin whiskers were found to have a much greater mechanical reinforcing effect than TiO2 (Fig. 4.5), owing to whisker/whisker interactions. The incorporation of tunicin whiskers induced an approximately threefold decrease in conductivity. Possible explanations include (i) the low dielectric constant of cellulosic fillers, (ii) interactions between cellulose and PEO, and (iii) the effect of whiskers on salt dissociation and ion mobility. Indeed, the relaxation time of 1H protons was significantly reduced with the addition of whiskers, and the diffusion coefficients for both cation and anion decreased by almost a factor of 3.120 The decrease in ionic mobility is in
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Log (E´)(Pa)
9 8 7 6 –100
–50
0 50 Temperature (°C)
100
150
4.5 Storage tensile modulus E′ vs. temperature at 1 Hz for the unfilled POE30–LiTFSI polymer electrolyte (䊉) and related composites filled with 10 wt% TiO2 (䊊) and tunicin whiskers (ⵧ).
good agreement with the reduction in ionic conduction. Thus the main effect seems to be the decrease in chain dynamics. Cationic transference numbers, T +, in the composites were calculated from the diffusion coefficient values. These values were almost constant for the composites, and were found to be temperature-independent. They ranged between 0.22 and 0.29 and are comparable to those reported for POE-based unfilled polymer electrolytes.125 Since the mechanical performance of the tunicin whisker/PEO–LiTFSI polymer electrolyte was ascribed to the percolating filler network, experiments were performed with plasticised polymer.124 PEO–LiTFSI-based polymer electrolyte with 30 wt% tetraethylene glycol dimethyl ether, TEGDME, as a plasticiser and tunicin whiskers as the nanometric reinforcing phase were processed. A thermal stabilisation effect above Tm was observed. The high temperature modulus, around 100 °C, was about 200 MPa for the composite PEO/TEGDME 30 wt% filled with 6 wt% whiskers and 500 MPa when filled with 10 wt% whiskers. This is an indication that plasticisation of the matrix did not interfere with the formation of the cellulosic network within the matrix. The conductivity of two different POE-based electrolytes was compared.124 It was observed that POE12–LiTFSI plasticised with 30 wt% TEGDME and filled with 6 wt% tunicin whiskers displayed a similar conductivity to the unplasticised and unfilled POE12–LiTFSI polymer electrolyte, but with much better mechanical properties.
4.6
Gel polymer electrolytes
A gel polymer electrolyte is a polymer swollen by a liquid electrolyte. These electrolytes exhibit high ionic conductivity at room temperature but present
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some reactivity with the lithium electrode owing to the presence of organic solvents; they also have poor mechanical properties. The challenges with these systems are therefore to improve their stability towards the lithium electrode, decrease gas evolution and enhance their mechanical properties. One solution might be to incorporate nanoparticles that would improve electrochemical stability and mechanical properties without decreasing ionic conductivity. The most common polymers used as the polymer matrix in gel polymer electrolytes are poly(vinylidene fluoride), PVDF, and polymethymethacrylate, PMMA. The nanoparticles used in gel polymer electrolytes are the same as those used in solid polymer electrolytes, i.e. organic clay126,127,128 and metal oxide fillers.129,130
4.6.1 Composite gel polymer electrolytes based on clays MMT is naturally hydrophilic and must be rendered hydrophobic or organophilic in order for it to be compatible with the polymer matrix. Gel nanocomposite polymer electrolytes were obtained by dispersing organophilic MMT in a PVDF or PVDF-hexafluoropropylene (PVDF-HFP) solution. The composite membranes were obtained by film-casting;128,131,132 the film was dried and then the liquid electrolyte was added by swelling.131,132 The MMT was exfoliated and, in order to improve MMT delamination, the MMT was sonicated in dimethylformamide (DMF).131 The filled PVDF membrane shown excellent wettability with the electrolyte solution. Solvent and electrolyte uptake by the filled membrane was considerably greater than that of the unfilled membrane.131 This increase in electrolyte uptake could explain the improvement in ionic conductivity observed. A similar study was performed by Jacob et al.126 using PVDF and PVDF– HFP as the polymer matrix, LiSO3CF3 as the salt and PC or EC as the solvent. A decrease in ionic conductivity was observed with the addition of clays, with conductivity equal to 4.4 × 10−4 S/cm for the filled gel electrolyte compared with 1.9 × 10−3 S/cm for the unfilled one, at room temperature. However, the conductivity of the filled electrolyte remained unchanged after several temperature cycles, while that of the unfilled electrolyte decreased owing to solvent evaporation. The nanoparticles may act as a physical and chemical barrier limiting solvent evaporation. The inconsistent results may be related to the difficulty in determining whether the added nanoparticles affect ionic conductivity directly, since the addition of nanoparticles can modify the amount of solvent and salt, which in turn could modify conductivity. The mechanical properties of nanogel electrolytes were also significantly improved. The nanogels exhibited a rubbery behaviour with a storage modulus of about 1 GPa at room temperature (one order of magnitude
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higher than the corresponding gels).126 MMT was incorporated in PMMA by dispersion in the monomer, and polymerisation was performed in situ.127,133 The gel PMMA–clay nanocomposites presented an intercalated structure and a slight increase in ionic conductivity was observed with 1.5 wt% clay.127,133 The PMMA nanocomposite gel electrolyte showed a stable lithium interfacial resistance, similar to PMMA gel. The mechanical properties of the pure PMMA were improved with the incorporation of 10 wt% MMT. The problem associated with this process is the need to purify the polymer by eliminating unreacted monomers.
4.6.2 Composite gel polymer electrolytes based on metal oxide fillers Nanocomposites based on PMMA were synthesised via sol–gel transformation and in situ free radical polymerisation of MMA. Gel nanocomposite polymer electrolytes were obtained by the addition of PMMA–TiO2 or PMMA–SiO2 to the PC-based liquid electrolyte. After 2 h at 55 °C, transparent films were obtained. The gel composite based on PMMA–TiO2 exhibited higher viscosity than the unfilled gel electrolyte and the PMMA– SiO2-based electrolyte. The conductivity values of filled and unfilled gel electrolytes were very similar, but the nanocomposites exhibited better mechanical properties than the unfilled polymer electrolytes. Nanocomposites based on PMMA were also obtained by SiO2 dispersion and PMMA solubilisation in PC/LiCF3SO3 liquid electrolyte.129 Improved ionic conductivity was obtained with 2 wt% SiO2 for temperatures higher than 50 °C, but at room temperature the addition had no effect on ionic conductivity. No change in Tg following incorporation of SiO2 was observed. The incorporation of fumed silica resulted in more physical linkages in the gel electrolyte in the form of siloxane linkages, which caused a substantial increase in the gel modulus. Wu et al.134 studied the effect of several nanocharges, TiO2, ZnO and MgO, on gel electrolytes based on PVDF–HFP. The gel electrolytes were obtained by mixing the nanocharges and the polymer in acetone. After solvent removal, the liquid electrolyte was added. A decrease in the porosity of the PVDF–HFP and an increase in solvent uptake were obtained with the addition of nanocharges. Owing to the higher amount of solvent introduced, an improvement in ionic conductivity was obtained. A PEO-based nanocomposite gel was studied by Nan et al.130 A solvent mixture EC/PC was added to mesoporous SiO2. The nanoparticles thus obtained were mixed with PEO–LiClO4 electrolyte. A decrease in PEO crystallisation was observed with the incorporation of EC/PC–SiO2. This decrease was greater than that observed with the addition of SiO2, which may be related to the presence of organic solvents. A net improvement in
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ionic conductivity was obtained with the incorporation of EC/PC–SiO2, even at high temperature. However, this increase may be related more to the incorporation of solvents than the direct effect of nanoparticles. The EC/PC–SiO2 filled PEO exhibited a very high Young’s modulus at room temperature, higher than the unfilled PEO electrolyte and EC/PC plasticised PEO electrolyte. This improvement in mechanical properties may be related to a specific interaction between SiO2 and PEO. Gel nanocomposite polymer electrolytes based on PVDF-HFP were studied by NMR.135 In this type of electrolyte, BaTiO3 and clay were used as nanoparticles and PC + LiCF3SO3 as the liquid electrolyte. The filled gel electrolytes presented better mechanical properties than the gels without the filler. The fillers lowered the conductivity by a small amount. The diffusion coefficient measured by NMR indicated that both anion and cation mobilities were reduced by the presence of the filler, but the effect was generally greater for the anions.
4.7
Composite electrolytes for proton exchange membrane fuel cells
Proton conducting membranes for proton exchange membrane fvel cells (PEMFCs) are usually based on thin ionomer films. The membrane separates the electrodes, allows proton transportation from anode to cathode and creates a barrier against the passage of gases or fluids (e.g. methanol). However, the ionomer itself does not provide any appreciable conductivity. The membrane must therefore be swollen by molecules, e.g. water, to ensure proton conductivity. Water uptake by the membrane is therefore one of the essential parameters for obtaining high conductivity. Indeed, owing to its high dielectric constant, water favours dissociation of the ion pair, increasing the concentration of charge carriers and the solvation of both anion and proton. Generally, two principal mechanisms are used to describe proton diffusion through the membranes.136–138 One is the vehicle mechanism, whereby a proton combines with vehicles such as H3O+ or CH3OH2+ and also with unprotonated vehicles (H2O), thus allowing the net transport of protons.136 In this mechanism, conductivity is directly dependent on the rate of vehicle diffusion. The other is the Grotthus mechanism (hopping), whereby protons are transferred from one proton acceptor site to another by hydrogen bonds. In this mechanism, additional reorganisation of the proton environment, consisting of reorientation of individual species or even more extended ensembles, leads to the formation of an uninterrupted trajectory for proton migration.136 Recently Marechal et al.139 found that four water molecules are involved in ion-pair solvation of Nafion® membranes. In Nafion®, which is the most advanced proton-conducting membrane, the drastic decrease in proton
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conductivity at low relative humidity limits the operating temperature to about 80 °C.140–149 An important focus for PEMFC research is developing a PEM (proton exchange membrane) which has satisfactory proton conductivity at higher temperatures (above 100 °C) and low relative humidity,140–142 because cells that can operate in such conditions have a number of advantages over lower temperature cells: •
Considerable improvement in the platinum tolerance of the electrocatalysts to carbon monoxide poisoning. Carbon monoxide is a major problem because trace amounts of CO in the H2 feed gas (more than 10 ppm) will poison the Pt anode electrocatalyst in PEMFCs operating at 80 °C. A quantitative analysis of the free energy for H2 and CO adsorption as a function of temperature150 suggests that, by elevating the operating temperature of the cell to 145 °C, CO tolerance at the anode should increase by a factor of 20 (from 5–10 to 100–200 ppm). • Significant improvement in PEMFC thermal management. By elevating the temperature of the fuel cell stack, thermal management can be simplified due to more efficient waste heat removal. • The amount of noble metal catalyst can potentially be reduced due to fast reaction kinetics at high temperature. • Water management is improved. At the moment, reactant gases need to be humidified before entering the cell to prevent the membrane from drying out. Membrane dehydration also causes the membrane to shrink, reducing contact between electrode and membrane and possibly leading to pinholes, which permit the crossover of reactant gases. To benefit from these advantages, it is now becoming clear that pure ionomer membranes are not suitable, and new types of proton-conducting membranes that work at temperatures higher than 100 °C have to be developed. There are currently three principal polymer electrolyte types proposed for high temperature PEMFCs: • •
•
Anhydrous proton-conducting membranes, in which proton conductivity is assured by molecules other than water. Hygroscopic inorganic/organic composite membranes, in which the ionomer or polymer is blended with hydroscopic inorganic particles so that the membranes will remain hydrated at higher operating temperatures. Self-humidifying membranes, in which the water is produced directly in the membrane.
Thermally stable anhydrous proton-conducting membranes involve membranes based on (i) acid–base composite membranes, i.e. acid blended with polymers containing Lewis base moieties, namely H3PO4/PBI (polybenz-
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imidazole),151–153 and (ii) proton-conducting ionic liquid composite membranes, i.e. a functional polymer blended with a proton-conducting ionic liquid.154–156 For the latter, an effective way to achieve low humidity and high temperature operation is to disperse particles of hygroscopic oxides or hygroscopic solid inorganic proton conductors with low solubility (e.g. heteropolyacids) inside the polymer matrix. In the case of self-humidifying membranes, research has focused mainly on incorporating Pt or Pt/C catalysts in the membrane to combine the permeable oxygen and hydrogen to produce water and humidify the membrane. These two types of polymer electrolytes are discussed below.
4.7.1 Components of composite electrolytes Polymer matrix Inorganic/Nafion® composites have been proposed for use at high temperatures in fuel cells.140–142,157–167 In addition to Nafion®, sulphonated hydrocarbon polymers such as (Scheme 4.3) sulphonated polystyrene (SPSt),168 sulphonated polysulphone (SPSF),169,170 sulphonated polyetheretherketone (SPEEK),171–175 and others without functional groups, such as PVDF,176–179 PEO180–183 and PBI,184–186 have been widely used as the host matrix for preparing inorganic/organic composites to be used at high operating temperatures. There are a number of reasons for the interest in developing inorganic– organic composite membranes: (i) to improve the self-humidification of the membrane on the anode side by dispersing a hydrophilic inorganic component homogeneously in the polymer; (ii) to reduce electro-osmotic drag and therefore prevent the membrane from drying out on the anode side; (iii) to eliminate fuel crossover, e.g. methanol in direct methanol fuel cell (DMFC); and (iv) to improve the mechanical strength of the membranes without sacrificing proton conductivity. In sulphonated polymers, for example, a high degree of sulphonation is desirable for high conductivity. However, this will be accompanied by undesirable swelling (or even solubility in water) of the membrane and therefore reduced mechanical strength. Introduction of an inorganic component into the polymer will compensate for the mechanical behaviour by helping to improve thermal stability and enhancing proton conductivity when solid inorganic proton conductors are used. This may permit the use of sulphonated polymers with a low degree of sulphonation. Inorganic particles Solid inorganic particles include oxides such as amorphous silica, inorganic proton conductors, and in many cases silica-supported inorganic proton
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PBI
N C
N
N N
H2 C
SPSt
CH2 n CH
CH
n
m
SO3H SPEEK
(
O
O
)n (O
CO
O
)m
CO
SO3H SPSF
O
R
O
SO2
n*
SO3H
O
R
O
SO2
p
m
p
R = –H or –C(CH3)2 SPAEEKK
O
O
O
O
CF3
C
C
C
70
O
O
O
C
C
30
CF3
SO3H SDFF
F
F F
F
CF3 O
F
F F
C O
O
CF3
F
n
SO3H SPPEK N N C=O
C
O n
O
SO3H
HO3S
Scheme 4.3 Polymer structures.
conductors. These compounds are crystalline and therefore mechanically poor when used alone in the form of membranes. The combination of these inorganic proton conductors with a polymer component provides flexibility. Composite polymer electrolytes The properties of these composite membranes depend not only on the nature of the ionomer and the solid used, but also on the amount, homo-
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geneous dispersion, size, and orientation of the solid particles dispersed in the polymer matrix.
4.7.2 Composite polymer electrolyte elaboration Composite membranes can be prepared according to two main procedures: • •
dispersion of inorganic particles in an ionomer solution followed by casting; and growth of the filler particles within a preformed membrane or in an ionomer solution.
The first attempts to disperse proton-conducting inorganic particles in an ionomer solution followed by film casting were performed over 25 years ago.187–188 In this procedure, the solids are first ground until a fine powder is obtained and then dispersed under strong stirring in an organic solution of the polymer. The membrane is obtained by film casting and solvent evaporation. The method is very simple, but it is usually difficult to prevent the formation of particle agglomerates inside the polymer matrix, even by the use of ultrasound, and thus membranes containing non-homogeneous dispersions of micro-sized particles are usually obtained. In order to reduce particle sizes and obtain more homogeneous membranes, Rosiere and co-workers175,189 developed colloidal dispersions of hydrous oxides. The advantage of colloidal solutions is that the solids are powdered into very small particles, which generally give homogeneous dispersions. Colloidal dispersions of hydrous oxides in aqueous solution are easily obtained by hydrolysis of metal alkoxides. Transfer of nanometresized silica particles from commercial aqueous silica suspensions to higher boiling point solvents such as N-methyl-2-pyrrolidinone (NMP) has been described.175,189 Such small and well-dispersed particles, preferentially localised in the hydrophilic domain of polymers, maximise the interface region between the inorganic and organic components. Mauritz and co-workers190 proposed a sol–gel approach involving the addition of silica to a Nafion® membrane, which might provide better performance than particles because the silica is homogeneously dispersed on the angstrom scale as opposed to submicrometre-length scales for particle composites. This method was developed for all types of ionomers. From a practical point of view, it is a very interesting way to modify commercial membranes. The filler precursors are incorporated in the polymer matrix by simple impregnation (if the filler is neutral) or by ion-exchange reaction (if the filler is a cationic species), and then the membrane is treated with the necessary reactants to convert the precursor into the final inorganic filler.191,192
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Another interesting family of PEMFCs is that of class II inorganic– organic hybrids, in which the inorganic part, generally silica, is covalently bonded to the polymer. These systems are more stable from a structural and mechanical point of view than other hybrids in which the inorganic units are only physically dispersed in the polymer matrix. Generally speaking, hybrid membranes have been prepared by sol–gel processes, starting with 3-isocyanatopropyltriethoxysilane and several polyethers. In this kind of system, the proton carriers are supplied by incorporating various types of acids into the membrane, in particular heteropolyacids, e.g. phosphotungstic (PWA) and silicotungstic acid (SWA) or H3PO4.180–183,193–196
4.8
Composite polymer electrolytes based on metal oxides
An effective way of achieving Nafion® membranes that operate under low humidity and high temperature conditions is to recast Nafion® membranes with mixed hygroscopic oxides (e.g. SiO2, TiO2, ZnO, Al2O3). Watanabe et al.197 presented the first report on this in 1994. It has been shown that water uptake by the oxide-containing membrane is higher than that of the pristine Nafion®. For recast Nafion® membranes pre-dried at 80 °C, the water absorbing ability under humidification with water vapour at 60 °C was found to be 17 wt%, whereas for membranes containing 3 wt% SiO2 of 7 nm size, the water absorbing ability was increased to as much as 43 wt%.197 The characteristics of composite membranes depend not only on the type of metal oxide, but also on surface area, size, concentration and surface chemical modification of the metal oxide, and on the type of polymer matrix. The primary method for characterising membranes is to compare the current–voltage response of the cell as a function of membrane chemical characteristics and cell humidity/temperature.
4.8.1 Electrochemical performance Antonucci et al.198 attributed the enhanced performance of the composite membrane to the hygroscopic properties of silica. Watanabe et al.197 investigated silica- and titania-impregnated Nafion® composite membranes in H2/O2 proton-exchange membrane fuel cells operating at 100 °C with minimal external humidification. They reported that the silica particles were superior to titania particles in terms of water retention qualities within the Nafion® membrane, and attributed this to the higher water sorption properties of silica. Adjemian et al.199,200 reported that the performances of silica and titania composites at 130 °C and 75% RH are comparable (0.33 Ω cm2 and 0.35 Ω cm2 cell resistivities) while recast Nafion®-based composites present high cell resistivity (0.9 Ω cm2).
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Some authors 146,199,200 have suggested that improved cell behaviour is not related to water retention or the acidity of the metal oxide component, but rather to the effect of the metal oxide particles on the temperaturedependent structure of the polymer matrix. To investigate the potential role of a chemical interaction between the metal oxide interface and the Nafion® sulphonate groups, Adjemian et al.199 constructed composite membranes containing alumina and zirconia particles. They reported that these systems behaved in a very similar manner, with both alumina and zirconia-based composites generally producing high resistance cells that were not tolerant of low humidity conditions. One possible reason given for the observed decrease in performance is the potential release of Al3+ or the production of zirconium oxysulphate in the acidic Nafion® environment, disturbing some of the sulphonate sites and leading to modification of the membrane topology, and thus potentially limiting proton diffusion.199 Park et al.201 compared the impact of physically mixed SiO2 + ZnO2 and binary SiO2–ZnO2 oxide fillers on Nafion® membrane properties. They concluded that the binary oxide improved proton conductivity and water uptake at high temperatures and low humidity. Cell performances improved at 120 °C with increasing silica content in the composite membrane containing the binary oxide. The binary oxide SiO2–ZnO2 (1–1) (10% w/w in Nafion®) gave the best performance. It is believed that the synergistic effects are heightened when both the oxides coexist in a single particle. The performance of binary oxide SiO2–ZnO2 (10%) was better than that of the physical mixture SiO2 + ZnO2 (10%) and higher than that of a membrane containing only one oxide in the same concentration (10%).
4.8.2 Interfacial characteristics Adjemian et al.199 examined the physicochemical characteristics of the various simple metal oxide composites, in particular the interfacial chemistry that occurs between the metal oxide particles and the Nafion® membrane. To investigate this, they processed titania samples using different treatments, degreasing or acidification in order to increase the surface hydroxyl content, and silylation of hydroxyl groups in order to reduce the surface hydroxyl content. In the first case, the treatments produced a membrane with improved PEMFC performance (resistivity value of 0.33 Ω cm2 and cell potential (at 600 mA/cm2) of 590 mV vs. 0.58 Ω cm2 and 582 mV for the untreated particles at 75% RH and 130 °C), suggesting surface contamination of the titania particles and that such surface contamination was detrimental to membrane performance. In the second case, the silylation led to an increase in resistivity (0.36 Ω cm2) and a decrease in cell potential (at 600 mA/cm2) to 400 mV under the same operating conditions, indicating
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Polymer electrolytes SO3H
Si-OH
Water molecules Proton
Proton transport
Scheme 4.4 Schematic representation of the model involved in proton-conduction mechanism and the state of water in the membrane of SPAEEKK/silica hybrid membrane.
that surface hydroxyl groups are the reactive sites needed to provide the requisite interaction between the titania particles and the Nafion® polymer. These results clearly show that the OH of the oxide surfaces are involved in the conduction mechanism. Mechanism Guiver and co-workers202 studied sulphonated poly (arylene ether ether ketone ketone) (SPAEEKK)/silica (Scheme 4.3) and proposed a mechanism for hybrid membranes (Scheme 4.4). They assumed that the nonfreezing and the freezing bound water participates by the Grotthus mechanism, and the free water participates by both vehicle and Grotthus mechanisms. The free water is evaporated with increasing temperature and consequently the contribution of the vehicle mechanism decreases. To maintain proton conductivity with increasing temperature, the decrease in the contribution of the vehicle mechanism must be compensated by an increase in the contribution of the Grotthus mechanism. Proton conductivity increased with silica content in all temperature conditions. This suggests that the silica doped into the membrane enhanced the formation of proton conduction pathways due to molecular water absorption. Thus, the probable mechanism for new proton conduction due to the presence of silica is proton hopping between Si-OH and water molecules.202,203 Nogami et al.203 reported that the dissociated proton moves to a water molecule bound with an SiOH bond, forming the activated H2O:H+ state
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(SiOH—H2O ⇒ SiO2 + H+:H2O). The proton from the activated H2O:H+ state dissociates to form a new activated state with a neighbouring H2O. They proposed that protons attached to protogenic surface functional groups of SiOH participate in the proton conduction process via the Grotthus mechanism.204
4.8.3 Physical/chemical characteristics of the matrix In addition to nanoparticle characteristics, the physical/chemical characteristics of the membranes have a strong impact on composite membrane properties. It was reported200 that the resistivity of the cell decreases as the equivalent weight and thickness of the Nafion® decreases in 10 wt% SiO2 Nafion® composites. For example, Nafion® 112 and 115 have the same equivalent weight of 1100 g of polymer per mol of sulphonic acid, while Nafion® 112 is 50 µm and Nafion® 115 is 125 µm thinner. The thinner Nafion® composite membrane (Nafion® 112) exhibits a lower resistivity (R = 0.53 Ω cm2 vs. 0.78 Ω cm2, 130 °C). Nafion® 105 and 115 have the same thickness of 125 µm, while Nafion® 105 has a higher sulphonate density (1000 equivalent weight versus 1100 equivalent weight). The lower equivalent weight Nafion® membrane (Nafion® 105) has a lower resistivity (R = 0.45 Ω cm2 vs. 1.3 Ω cm2). At a cell temperature of 130 °C, Nafion® 105 exhibits the lowest resistivity of the extruded Nafion® membranes, indicating that equivalent weight plays a larger role than membrane thickness in maintaining proton conductivity at 130 °C.
4.8.4 Metal oxide concentration Most publications193–196 reported that proton conductivity increased with metal oxide content under all temperature conditions. This suggests that the metal oxide doped into the membrane improves the formation of proton conduction pathways due to molecular water absorption. However, some authors observed146,205 an improvement in the mechanical properties of polymer electrolytes with the addition of nanofillers at optimal concentrations. The optimal concentration depended not only on the polymer and inorganic oxide structure, but also on the elaboration procedure. Adjemian et al.146 showed that 6% SiO2 by weight was the optimal percentage of silicon oxide for extruded Nafion®, while 10% by weight was optimal for recast Nafion®. An excess of the metal oxide phase causes not only low proton conductivity but also poor physical properties, which suggests that an excess of the metal oxide phase prevents the organic–inorganic hybrid membranes from forming an interpenetrated network. A composite membrane with sulphonated poly(decafluorobiphenyl-4-4′(hexafluoroisopropylidene) diphenol) (SDFF)206 (Scheme 4.3) + 4 wt% by
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Tensile strength (MPa)
30 25 20
∗ Nafion
15 10 5 0
0
2
4 6 8 SiO2 content (%)
10
12
4.6 Tensile strength of membranes as a function of SiO2 content: (䊏) polymer membrane and composite membrane; (*) Nafion® film, adapted from Kim et al.206
Table 4.1 Proton conductivity of membrane containing various SiO2
Conductivity (S/cm)
SDFF
SDFF–SiO2 (4%)
SDFF–SiO2 (7%)
SDFF–SiO2 (10%)
SDFF–SiO2 (18%)
1.05 × 10−5
3.9 × 10−5
3.7 × 10−5
3 × 10−5
Very brittle
weight SiO2 shows significant improvement, especially at high temperature. A higher silica concentration reduces conductivity and mechanical properties (Fig. 4.6, Table 4.1).
4.9
Hygroscopic solid inorganic proton conductor composite polymer electrolytes
4.9.1 Hygroscopic solid inorganic proton conductor Hygroscopic solid inorganic proton conductors as membrane fillers increase the number of protonic carriers and thus improve the hydrophilic character of the membranes. Among the inorganic solid proton conductors,137,200,207,208 zirconium phosphates, heteropolyacids, metal hydrogen sulphate and a few others are of special interest for their potential role in developing high temperature composite membranes for PEMFCs.
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Zirconium phosphates (ZP) have long been recognised as interesting inorganic ion exchangers,209 especially after their crystalline form Zr(HPO4)2 · 2H2O was prepared in the 1960s.210 This group of compounds can be expressed as MIV(RXO3) · nS, where M is a tetravalent metal such as Zr, Ti, Ce, Th or Sn; R is an inorganic or organic group such as –H, –OH, –CH3OH or (CH2)n; X is P or As; and S is a solvent, i.e. H2O. They form two types of layered structures, known as α and γ.137,211 These compounds exhibit good proton conductivity in a temperature range up to 300 °C. In the form of glassy plates or films, a conductivity of 10−2 S/cm at room temperature has been reported.212 The conductivity of Zr(HPO4) · 2H2O and H3O–Zr2(PO4)3 at 100 °C and 100% RH was determined to be about 10−1 S/cm213,214 and 7 × 10−7 S/cm,215 respectively. Heteropolyacids (HPA) possess strong acidity and high proton conductivity in their hydrated form,139 and their molecules (about 1 nm in diameter) can be regarded as nanoparticles. HPAs typically exist in hydrated phases, with the degree of hydration varying from 30 to 6 molecules of water (waters of hydration) per HPA molecule.216–218 The proton conductivities of these molecules drops respectively from approximately 0.06 to 2 × 10−5 S/ cm.214 The exact number of waters of hydration depends on the temperature and relative humidity of the environment. At 120 °C and 35% relative humidity, 12-phosphotungstic acid ((H3PW12O40 · nH2O) PWA) has approximately 6 waters of hydration.214 Detailed information about various HPA structures can be found in the literature.218–221
4.9.2 Composite polymer electrolytes based on zirconium phosphate Chalkova et al.222 studied the effects of ZP particle size, structure and distribution in a membrane on the performance of Nafion®/ZP composite membranes in PEMFCs at high temperatures and low RH. They synthesised Nafion®/ZP composite membranes with different structures containing in situ precipitated ZP, crystalline layer-structured α-zirconium hydrogen phosphate (Zr(HPO4) · 2H2O) (α-ZPL), and three-dimensional-network zirconium hydrogen phosphate H3O–Zr2(PO4)3 (ZPTD). The difference in performance of composite Nafion®/ZP membranes containing ZP in different structures was more pronounced at 120 °C than at 80 °C, particularly at 13 and 26% RH. The Nafion®/α-ZPL composite membranes demonstrated the best performance for all RH values. The advantage of these membranes over Nafion®/ZPTD and Nafion®/ZP in situ membranes was especially obvious at 13% RH. The current density obtained at a cell voltage of 0.6 V for the Nafion®/α-ZPL membrane at 13% RH was 10 times higher than that of an unmodified recast Nafion® membrane at 26% RH. Membranes with 10%
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in situ precipitated ZP demonstrated much lower performance in a PEMFC than membranes with 10% α-ZPL or ZPTD. To enhance the performance of the Nafion®/ZP in situ precipitated composite membrane, the concentration of ZP nanoparticles has to be substantially increased. Improved performance of Nafion®/ZP membranes in a fuel cell were reported for membranes containing 23 and 36 mass% of ZP.219,223 These studies suggested that, first, the effect of pore constriction219,222 and/or, second, scaffold formation222,224 were determining factors for improving membrane performance. In the former case, it was possible that 10% ZP was not enough to fill all the available pores and induce a capillary condensation effect, which enhances water retention in the membrane. In the latter, 10% ZP could also be insufficient to form internal rigid scaffolding within the whole membrane, which would help to resist compression due to the sealing pressure in a fuel cell and facilitate water uptake by the membrane at a reduced RH.219 The difference in performance of Nafion®/ZP composite membranes containing α-ZPL and ZPTD can be attributed to the structure and surface properties of the inorganic particles and membrane morphology. Different structural modifications of ZP have different types of proton diffusion path, which affect proton conductivity through the inorganic material framework.225 In layered structures, the bulk diffusion of the hydronium ions occurs through the open cavities parallel to the layer,226 but the mobility of interior hydronium ions is about 104 times less than the mobility of surface (adsorbed) hydronium ions227 because the free water molecules on a crystal surface can easily rotate and contribute to ionic conduction by the Grotthus mechanism. It is difficult for hydronium ions to rotate in the interlayer spaces of ZP crystals due to strong hydrogen bonding with layered phosphate groups. As a result, proton transport in α-ZPL is dominated by surface transport.222,228 In a three-dimensional structure, conduction at room temperature is believed to occur via the Grotthus mechanism, mainly through grain boundaries. As the temperature increases, the surface water is lost, and grain boundary conduction decreases, being replaced by classical H+-ion hopping through the cavities inside the crystal structure.229 Particle distribution in composite membranes can also be regarded as a factor affecting membrane performance in a fuel cell. Chalkova et al.222 reported improved performance for α-ZPL, explained by the negative zeta potential of this filler, which provides electrostatic dispersion of particles in the acidic suspension during membrane casting and prevents aggregation and settling. The fact that the zeta potential of α-ZPL was more negative than that of ZPTD suggests that its surface phosphate groups undergo more complete dissociation (–POH = −PO− + H+) in an acidic environment. Composite Nafion®–Teflon®-α-ZPL membranes, prepared by incorporating Nafion® and α-ZPL into the Teflon® matrix, have been developed
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and tested in an H2/O2 fuel cell at 120 °C, ambient pressure and 35% RH by Si et al.230 Improved performance was due to both the incorporation of α-ZPL filler and reduced membrane thickness. Hydrogen crossover for the composite membrane was about 2.5 times higher than for a commercial Nafion® 1135 membrane (2 mA/cm2 with Nafion®-Teflon®-α-ZPL vs. 0.8 mA/cm2 Nafion® 1135) at 120 °C, which is still low enough not to cause significant performance loss. Kim et al.231 prepared Nafion®/zirconium sulphophenyl phosphate (ZSPP) composite membranes. ZSPP was precipitated by the reaction of Zr4+ ion and m-sulphophenyl phosphonic acid (SPP) with a stoichiometric ratio P/Zr = 2. They found the conductivity of ZSPP to be about 10−2 S/cm, much higher than ZPL.216 The conductivity increased with the crystallinity of ZSPP. The increased conductivity may be related to more conduction channels through the interlayer of crystalline ZSPP. When 12.5 wt% ZSPP particles was incorporated in Nafion®, it showed quite high conductivity, 0.07 S/cm at 140 °C.
4.9.3 Composite polymer electrolytes based on heteropolyacids Nafion® recast membranes filled with silica-immobilised phosphotungstic acid, PWA/SiO2 (30/70 weight ratio) and silicotungstic acid, SWA/SiO2 (45/65 weight ratio), so that the silica content was in all cases 3 wt%, were prepared by bulk mixing a 5% Nafion® solution with the powdered filler and then thermally treating it at 160 °C.232 Tests in direct methanol fuel cells at 145 °C showed that the PWA/SiO2 membrane had better electrochemical characteristics at high current densities compared with membranes loaded with SWA/SiO2 or only 3 wt% SiO2. Peaks of power density of 400 and 250 mW/cm2 were reached by feeding the cell with oxygen and air. For the same fillers, Tian and Savadogo233 reported better improvements for Nafion® +PWA/SiO2 membranes obtained by casting from DMF (0.3 A/ cm2 at 0.6 V) solution than for those obtained by aqueous solution (0.15 A/ cm2 at 0.6 V). Hongwei et al.234 found that sulphonated poly(phthalazinone ether ketone) (SPPEK)/12-PWA composite membranes showed improved conducting properties, and the SPPEK/ZrP composite membrane showed similar conducting properties to the pure SPPEK membrane. The SPPEK/ PWA composite membrane showed the highest proton conductivity, 0.17 S/ cm at 80 °C. The difference in conducting properties between the SPPEK/ PWA composite membrane and the SPPEK/ZrP composite membrane can be explained by their conduction mechanism. Kim et al.235 obtained hybrid PEMs based on polytetramethylene oxide (PTMO) and polydimethysiloxane (PDMS) using the sol–gel technique.
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The organic–inorganic interfaces were formed by sol–gel processing the end groups (OH) of organic parts with reactive inorganic parts, leading to hybrids at the molecular scale. The PWA was incorporated using sol–gel processes and the effects of PWA doping concentration on the thermal stability and protonic conductivity were discussed.236–239 The hybrid membranes showed superior thermal and mechanical properties because of the temperature-tolerant organic–inorganic moieties and the nanostructure of the hybrid matrix. The maximum proton conductivities for zirconium phosphate–PTMO membranes and titania–PTMO–40 wt% PWA hybrids were 2 × 10−3 and 8 × 10−2 S/cm, respectively, under saturated humidity conditions such that PWA or phosphate are bonded with hydrophilic metal phases. The cell performance of the membrane electrode assembly (MEA) with zirconium phosphate–PTMO membrane showed a maximum power density of 13 mW/cm2 and the titania–PTMO–PWA composite membrane with Nafion® supporter showed a maximum power density of 30 mW/cm2. Moreover, the organic–inorganic hybrid membrane showed stable performance up to 130 °C under saturated humidity conditions. Smitha et al.240 prepared solution cast films of composite membranes containing 40 wt% PWA and 60 wt% polysulphone (PSF) and SPSF. The PWA–40/SPSF composite membranes showed an increase in Tg, suggesting an intermolecular interaction between PWA and SPSF, while a reduction in the Tg of the PWA–40/PSF composites containing 40% PWA was found, indicating the absence of physicochemical interactions. Fourier transform infrared (FTIR) band shifts of this composite showed that sulphonic acid groups interact with both bridging tungstic oxide and terminal tungstic oxide, implying that the hydrogen bonding is attributed solely to miscibility and physicochemical properties. Incorporation of PWA into the sulphonated polymer significantly reduced the water swelling behaviour of the composite membrane. The membranes exhibited high conductivity (0.09 S/ cm) at room temperature, and a maximum of about 0.14 S/cm at 120 °C. Another sulphonated polysulphone composite membrane, obtained by bulk mixing of powdered H3Sb3P2O14 with the polymer solution in dichloroethane-isopropanol, exhibited higher conductivity, lower water swelling and lower gas permeability than unmodified SPSF. The conductivity of these membranes showed a maximum for 8 wt% filler. For this loading, the conductivity of SPSF with an ion exchange capacity of 1.1 mequivalent/g was enhanced by a factor of 10 and reached 2 × 10−2 S/cm at 80 °C and 100% RH.169,170
4.10
Self-humidifying composite electrolytes
Another class of inorganic/organic composite membranes, called ‘selfhumidifying’ polymer electrolyte membranes, have been developed.241–244
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These membranes contain highly dispersed nanosized Pt, which stimulates catalytic recombination of H2 and O2 inside the membranes, and a hygroscopic metal oxide, which adsorbs the water produced at the Pt particles and at the cathode. Self-humidifying membranes, in particular Nafion®–Pt– SiO2, Nafion®–Pt–TiO2 and Nafion®–Pt–zeolite,245,246 exhibited significantly higher water uptake and lower fuel crossover than unmodified Nafion® membranes. However, it is difficult to have uniformly distributed Pt particles in hygroscopic materials made by conventional methods such as electroless plating, sputtering or Pt black doping. There is also the possibility of a short circuit forming in the membrane because of non-uniformly dispersed Pt particles. In order to avoid a short circuit after incorporating Pt or Pt/C particles, another type of self-humidifying membrane has recently been developed. This is a thin double-layer composite membrane consisting of one layer of Pt/C catalyst-dispersed Nafion® and another layer of recast Nafion®. Researchers have also developed three-layered membrane structures to resolve this problem.247–249 Yang et al.249 reported that, with dry H2 and O2, the composite membrane showed about 90% of the performance obtained with the humidified reactant under similar conditions. However, the fabrication processes for these membranes are too complex for widespread use. To improve PEMFCs with dry reactants, an inorganic/organic self-humidifying membrane based on a SPEEK hybrid with Cs2.5H0.5PW12O40 supported Pt catalyst (Pt-Cs2.5 catalyst) was investigated.250 The Pt-Cs2.5 catalysts incorporated in the SPEEK matrix provide sites for catalytic recombination of permeable H2 and O2 to form water, and meanwhile avoid short circuits through the whole membrane thanks to the insulating properties of the Cs2.5H0.5PW12O40 support. Furthermore, the Pt-Cs2.5 catalyst can adsorb the water and transfer the proton inside the membrane. The physicochemical and electrochemical properties, such as ion exchange capacity (IEC), water uptake and proton conductivity, between the plain SPEEK and the SPEEK/ Pt-Cs2.5 composite membrane were compared. Additive stability measurements indicated that the Pt-Cs2.5 catalyst showed improved stability in the SPEEK matrix compared to the PWA particles in the SPEEK matrix. Single cell tests employing the SPEEK/Pt-Cs2.5 self-humidifying membrane and the plain SPEEK membrane under wet or dry operating conditions and primary 100 H fuel cell stability measurements were also conducted by these authors. The performance of cells using self-humidifying composite membranes has improved considerably compared with that of composite membranes without self-assembled Pt nanoparticles. However, the performance of selfhumidifying composite membranes without external humidification is still inferior to that of cells with external humidification. At a current density of 900 mA/cm2, cell potential with external humidification at 80 °C was found
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to be 0.603 V, higher than for the cell without external humidification (0.402 V). Similar improvements in cell performance were also observed at an operating temperature of 60 °C. This indicates that the self-humidifying effect of the self-assembled Pt nanoparticle layer may not be sufficient to humidify the membrane under the conditions investigated in the literature. As a result, the advantages of self-humidifying PEMs lie in their ability to reduce the need for external humidification when excellent performance is required.
4.11
Future trends
For polymer electrolytes based on polyether used in lithium batteries, it is clear that the addition of nanofillers has a positive impact on ionic conductivity at temperatures below the membrane composite melting temperature. At temperatures above the membrane melting temperature, the improvement in ionic conductivity is questionable. The use of nanocharges with high dielectric constant or/and which induce specific interactions with the anion might be worth investigating. Basic research is needed to confirm the conduction mechanism in nanocomposite electrolytes. When studying the effects of fillers, the nanofillers themselves and the composite electrolytes must be dried carefully to ensure that any observed changes are not due to the effects of water trapped by the fillers. Mechanical improvement can be obtained by incorporating large aspect ratio fillers, which may be an advantage if they allow the thickness of the electrolyte to be reduced, thus lowering internal resistance. However, more research is needed in this field in order to get a good compromise between mechanical properties and ionic conductivity. In gel polymer electrolytes, the addition of fillers can have beneficial effects on both liquid electrolyte uptake and mechanical properties. However, as for SPEs, composite gel polymer electrolytes must be performed carefully in order to remove water trapped in the fillers. The current challenge for PEMFC is to raise the working temperature above 80 °C. Composite membranes are a potential solution. The addition of inorganic fillers induces important improvements in water retention at high temperature, conductivity, cell resistivity, mechanical properties, etc. These improvements are related to filler concentration, structure and size, interfaces, polymer matrix and membrane characteristics. It is difficult to compare these ionomer/filler composites because their performances depend on the electrolyte preparation and testing conditions (RH, temperature, etc.). H2/O2(air) cells based on composite polymer electrolytes have been successfully operated at temperatures up to 120 °C under ambient pressure, and up to 150 °C under pressures of 3–5 atm, but more research
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is needed into their application in high temperature PEMFC in order to establish how long these composite electrolytes continue to exhibit high performance at low relative humidity and high temperature. An interesting idea is to produce water within a membrane by catalytic recombination of H2 and O2 on a platinum catalyst dispersed in the membrane. One of the difficulties in developing such electrolytes is preventing short circuits, which requires a uniform dispersion of Pt inside the membrane, another is the prohibitive price. Since the water produced within the membrane is insufficient to ensure good performance, external humidification is essential. The advantage of self-humidifying PEMs lies in their ability to reduce the humidity of the reactant gases. Anhydrous proton conducting electrolytes, such as electrolytes based on proton conducting ionic liquids, might be better adapted for PEMFC working at temperatures higher than 120 °C.
4.12
Sources of further information and advice
Books on composite electrolytes and polymer composites can be consulted for complementary information. Indeed, the properties of composite materials are associated with both components, the matrix and the charge, and also with interfacial properties. Some interesting data and information on experimental investigations in the composite polymer field can be found in the following books: •
•
•
•
•
Hydrocarbon Polymer Electrolytes for Fuel Cell Application (J. Qiao and T. Okada, 2008, Nova Science Publishers) focuses on composite membranes for direct methanol fuel cells. Polymer Electrolyte Fuel Cell Durability (F.N. Büchi, M. Inaba, and T.J. Schmidt, (Eds), 2009, Springer) covers one of the most important aspects of fuel cells, durability, and gives pertinent information on polymer membrane degradation, which may be an important orientation for future studies on composite membranes. Composite Polymeric Electrolytes (W. Wieczorek and M. Siekierski, 2007, Springer) reviews the considerable experience of W. Wieczorek and his team in work on composite electrolytes for lithium batteries. Polymer Composite Materials Interface Phenomena and Processes (Y. Ivanov, V. Cheshkov and M. Natova, 2001, Kluwer Academic) evaluates the properties of polymer composites in relation to acid–base properties. Handbook of Polymer Blends and Composites (C. Vasilen and A.K. Kulshrehtha, 2003, Woodhead Publishing) provides an overwiew of the theory and pratice of polymer blends and composites.
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170. baradie, b., poinsignon, c., sanchez, j.y., piffard, y., vitter, g., bestaoui, n., foscallo, d., denoyelle, a., delabouglise, d., vaujany, m. 1998, J. Power Sources, 74, 8. 171. bonnet, b., jones, d.j., roziere, j., tchicaya, l., alberti, g., casciola, m., massinelli, l., baner, b., peraio, a., ramunni, e. 2000, J. New Mat. Electrochem. Syst., 3, 87. 172. zaidi, s.m.j., mikhailenko, s.d., robertson, g.p., guiver, m.d., kaliaguine, s. 2000, J. Membr. Sci., 173, 17. 173. nunes, s.p., ruffmann, b., rikowski, e., vetter, s., richau, k. 2002, J. Membr. Sci., 203, 215. 174. mikhailenko, s.d., zaidi, s.m.j., kaliaguine, s. 2001, Catal. Today, 67, 225. 175. roziere, j., jones, d.j., tchicaya-bouckary, l., bauer, b. 2000, Int. Patent No. WO 02/05370. 176. navarra, m.a., fernicola, a., panero, s., scrosati, b. 2006, J. Electrochem. Soc., 153, A1284. 177. croce, f., hassoun, j., tizzoni, c., scrosati, b. 2006, Electrochem. Commun., 8, 1125. 178. peled, e., duvdevani, t., melman, a. 1998, Electrochem. Solid-State Lett., 1, 210. 179. shen, j., xi, j., zhu, w., chen, l., qiu, x. 2006, J. Power Sources, 159, 894. 180. honma, i., takeda, y., bae, j.m. 1999, Solid State Ionics, 120, 255. 181. nakajima, h., honma, i. 2002, Solid State Ionics, 148, 607. 182. honma, i., nomura, s., nakajima, h. 2001, J. Membr. Sci., 185, 83. 183. nakajima, h., nomura, s., sugimoto, t., nishikawa, s., honma, i. 2002, J. Electrochem. Soc., 149, A953. 184. staiti, p., minutoli, m., hocevar, s. 2000, J. Power Sources, 90, 231. 185. staiti, p., minutoli, m. 2001, J. Power Sources, 94, 9. 186. staiti, p. 2001, Mater. Lett., 47, 241. 187. alberti, g., casciola, m., costantino, u., levi, g. 1978, J. Membr. Sci., 3, 179. 188. alberti, g., casciola, m., costantino u. 1983, J. Membr. Sci., 16, 137. 189. tchicaya-bouckary, l., jones, d.j., roziere, j. 2002, Fuel Cell, 2, 40. 190. deng, q., moore, r.b., mauritz, k.a. 1995, Chem. Mater., 7, 12, 2259. 191. grot, w., rajendran, g. 1996. Int. Patent No. WO 96/29752. 192. mauritz, k.a., warren, r.m. 1989, Macromolecules, 22, 1730. 193. nakamura, o., kodama, t., ogino, i., miyake, y. 1979, Chem. Lett., 1, 17. 194. honma, i., nakajima, h., nomura, s. 2002, Solid State Ionics, 154–155, 707. 195. nakajima, h., nomura, s., sugimoto, t., nishikawa, s., honma, i. 2002, J. Electrochem. Soc., 149, A953. 196. lavrencic, s., stangar, u., orel, b., vince, j., jovanovski, v., spreizer, h., surca vuk, a., hocevar, s. 2005, J. Solid State Electrochem., 9, 106. 197. watanabe, m., uchida, h., seki, y., emori paul stonehart m. 1996, J. Electrochem. Soc., 143, 3847. 198. antonucci, p.l., arico, a.s., creti, p., ramunni, e., antonucci, v. 1999, Solid State Ionics, 125, 431. 199. adjemian, k.t., krishnan, l.d., ota, h., majsztrik, p.t., zhang, p., mann, j., kirby, b., gatto, l., velo-simpson, m., leahy, j., srinivasan, s., benziger, j.b., bocarsly, a.b. 2006, Chem. Mater., 18, 2238. 200. adjemian, k.t., srinivasan, s., benziger, j., bocarsly, a.b. 2002, J. Power Sources, 109, 356.
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201. park, k.t., jung, u.h., choi, d.w., chun, k., lee, h.m., kim, s.h. 2008, J. Power Sources, 177, 247. 202. kim, d.s., michael, b.l., guiver, d. 2006, Polymer, 47, 7871. 203. nogami, m., nagao, r., wong, c. 1998, J. Phys. Chem. B, 102, 5772. 204. krishnan, p., park, j.s., kim, c.s. 2006, J. Membr. Sci., 279, 220. 205. mustarelli, p., carollo, a., grandi, s., quartarone, e., tomasi, c., leonardo, s., magistris, a. 2007, Fuel Cells, 6, 441. 206. kim, y.m., choi, s.h., lee, h.c., hong, m.z., kim, k., lee, h-i. 2004, Electrochim. Acta, 49, 4787. 207. reyna-valencia, a., kaliaguine, s., bousmina, m. 2005, J. Appl. Polym. Sci., 98, 2380. 208. alberti, g., casciola, m. 2001, Solid State Ionics, 145, 3. 209. alberti, g., casciola, m., costantino, u., vivani, r. 1996, Adv. Mater., 8, 291. 210. clearfield, a., stynes, j.a. 1964, J. Inorg. Nucl. Chem., 26, 117. 211. alberti, g., casciola, m. 1997, Solid State Ionics, 97, 177. 212. abe, y., li, g., nogami, m., kasuga, t., hench, l.l. 1996, J. Electrochem. Soc., 143, 144. 213. arico, a.s., baglio, v., di blasi, a., antonucci, v. 2003, Electrochem. Commun., 5, 862. 214. hogarth, w.h.j., diniz da costa, j.c., lu, g.q. 2005, J. Power Sources, 142, 223. 215. subramanian, m.a., roberts, b.d., clearfield, a. 1984, Mater. Res. Bull., 19, 1471. 216. ramani, v., kunz, h.r., fenton, j.m. 2004, J. Membr. Sci., 232, 31. 217. colomban p. (Ed.) 1992, Chemistry of Solid State Materials. 2. Proton Conductors – Solids, Membranes and Gels – Materials and Devices, Cambridge University Press, Cambridge. 218. nakamura, o., ogino, i. , kodama, t. 1981, Solid State Ionics, 3–4, 347. 219. pope, m.t. 1983, Heteropoly and Isopoly Oxometalates, Springer-Verlag, Berlin. 220. moffat, j.b. 1986, Polyhedron, 5, 261. 221. brown, g.m., noe-spirlet, m.r., busing, w.r., levy, h.a. 1977, Acta Cryst., B33, 1038. 222. chalkova, e., fedkin, m.v., komarneni, s., lvova, s.n. 2007, J. Electrochem. Soc., 154, B288. 223. yang, c., srinivasan, s., arico, a.s., creti, p., baglio, v., antonucci, v. 2001, Electrochem. Solid-State Lett., 4, A31. 224. yang, c., srinivasan, s., bocarsly, a.b., tulyani, s., benziger, j.b. 2004, J. Membr. Sci., 237, 145. 225. clearfield, a., roberts, b.d., subramanian, m.a. 1984, Mater. Res. Bull., 19, 219. 226. ticianelli, a., derouin, c.r., redondo, a., srinivasan, s. 1998, J. Electrochem. Soc., 135, 2209. 227. alberti, g., casciola, m., costantino, u., levi, g., ricciardi, g. 1978, J. Inorg. Nucl. Chem., 40, 533. 228. casciola m., bianchi, d. 1985, Solid State Ionics, 17, 287. 229. trakanprapai, c., esposito, v., licoccia, s., traversa, e., baglio, v., di blasi, a, arico, a.s., antonucci, v., antonucci, p.l. 2003, Abstract 1080, The Electrochemical Society Meeting Abstracts, Vol. 2003-2, Orlando, FL, Oct. 12. 230. si, y., kunz, h.r., fenton, j.m. 2004, J. Electrochem. Soc., 151, A623. 231. kim, y-t., song, m-k., kim, k-h., park, s-b., min, s-k., rhee, h-w. 2004, Electrochim. Acta, 50, 645.
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232. staiti, p., aricio, a.s., baglio,v., lufrano, f., passalacqua, e., antonucci, v. 2001, Solid State Ionics, 145, 101. 233. tian, h., savadogo, o. 2005, Fuel Cell, 5, 375. 234. hongwei, z., baoku, z., youyi, x. 2006, J. Appl. Polym. Sci., 102, 3972. 235. kim, j.-d., mori, t., honma, i. 2006, J. Electrochem. Soc., 153, A508. 236. coronado, e., garcia, c.j.g. 1998, Chem. Rev., 98, 273. 237. nakamura, o., kodama, t., ogino, i., miyake, y. 1979, Chem. Lett., 17. 238. honma, i., nakajima, h., nomura, s. 2002, Solid State Ionics, 154/155, 707. 239. honma, i., nakajima, h., nishikawa, o., sugimoto, t., nomura, s. 2002, J. Electrochem.Soc., 150, 616. 240. smitha, b., sridhar, s., khan, a.a. 2005, J. Polym. Sci.: Part B: Poly. Phys., 43, 1538. 241. watanabe, m., uchida, h., seki, y., emori, m. 1996, J. Electrochem. Soc., 143, 3847. 242. yang, t.h., yoon, y.g., kim, c.s., kwak, s.h., yoon, k.h. 2002, J. Power Sources, 106, 328. 243. kwak, s.h., yang, t.h., kim, c.s., yoon, k.h. 2003, J. Power Sources, 118, 200. 244. uchida, h., ueno, y., hagihara, h., watanabe, m. 2003, J. Electrochem. Soc., 150, A57. 245. wang, l., xing, d.m., liu, y.h., cai, y.h., shao, z.g., zhai, y.f., zhong, h.x., yi, b.l., zhang, h.m. 2006, J. Power Sources, 161, 61. 246. son, d.h., sharma, r.k., shul, y.g., kim, h.s. 2007, J. Power Sources, 165, 733. 247. zhu, x., zhang, h., liang, y., zhang, y. 2006, Electrochem. Solid-State Lett., 9, A49. 248. zhang, y., zhang, h.m., zhu, x.b., liu, g., bi, c., liang, y.m. 2007, J. Power Sources, 165, 786. 249. yang, y.z., fu, a., manthiram, r. 2005, J. Power Sources, 139, 170. 250. zhang, y., zhang, h., bi, c., zhu, x. 2008, Electrochim. Acta, 53, 4096.
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5 Lithium-doped hybrid polymer electrolytes V. de ZEA B E R M U D E Z, University of Trás-os-Montes e Alto Douro, Portugal, and M. M. S I LVA, University of Minho, Portugal
Abstract: The interest in sol–gel derived polymer/siloxane hybrid electrolytes has grown considerably during the last decade because of their potential application in solid state electrochemical devices, in particular batteries and electrochromic devices (ECDs). This review intends to provide an overview of the latest advances in the investigation of the structure, morphology, thermal properties, electrochemical behavior and spectroscopic features of Li+-containing di-urea crosslinked polyoxyethylene (POE)/siloxane (di-ureasil) electrolytes. Applications of these materials in ECDs will be addressed. Key words: hybrid electrolytes, sol–gel, lithium salts, characterization, electrochromic devices.
5.1
Introduction
The interest in polymer electrolytes (PEs) has not ceased to increase since the first report of salt/polymer complexes by Wright1 and the subsequent recognition by Armand et al.2 of the potential use of these materials as electrolytes in solid state electrochemical devices.3 Highly conducting, thin, solvent-free PE films may be produced simply by dissolving ionic salts in high molecular weight polymers carrying appropriate coordinating species, such as ether oxygen atoms. Owing to its extraordinary solvating ability towards salts, polyoxyethylene (POE) has been the archetypal host polyether employed to synthesize PEs. Within this category of electrolytes, those incorporating Li+ ions have been the most extensively investigated to date.4 These materials fulfil the prerequisites of environmental compatibility, offer no hazard to the consumer, are of low density and can be produced as thin films. They are particularly suitable for application in advanced batteries (primary and secondary) and ECDs. The practical application of POE-based electrolytes in solid state electrochemical devices has been delayed by two main drawbacks: their poor processability and marked tendency to crystallize. The latter aspect is particularly limiting. POE/salt complexes are usually crystalline, but pristine POE is itself already semicrystalline. For nearly two decades the view that 176 © Woodhead Publishing Limited, 2010
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prevailed in the PE community, based on the pioneer work of Berthier et al.,5 was that in semicrystalline PEs ionic transport was confined to amorphous phases. It was accepted that high ionic conductivity could be expected in systems composed of host polymers with low glass transition temperatures (Tg). Understandably, until very recently the elucidation of the phase diagram of semicrystalline electrolytes was considered to be indispensable to determine the ranges of salt concentration and temperature for which the conductivity attained a maximum value. In 2001 Gadjourova et al.6 showed for the first time that ionic conductivity in the crystalline domains of the polymer matrix can be significantly higher than in the corresponding amorphous regions. However, because of the low conductivity exhibited by the complexes proposed by these authors6 current research remains focused on the development of entirely amorphous systems. To overcome the disadvantages associated with conventional POE-based PEs, three strategies have been basically attempted in the last two decades: the modification of the polymer architecture, the modification of the anion design and the addition of plasticizers.3 The class of organically modified silicate electrolytes (ormolytes) that emerged from these efforts has attracted considerable attention. The ormolyte strategy relies on the production in mild synthetic conditions of organic/ inorganic hybrid host frameworks7 by means of the sol–gel method.8 The incorporation of POE segments in the organic component endows these siloxane-based hybrid structures with the ability to accommodate ionic species. Interestingly, in these hybrid electrolytes the proportion of crystalline regions is in general very low or even null. Moreover, extremely high contents of salt may be added to the hybrid matrices, while avoiding the inconvenient ‘salting-out’ often found in classical PEs. Xerogel samples of these easy-to-process materials may be readily produced as thin films and a remarkable improvement in the mechanical properties and chemical/ thermal stability results. The organic/inorganic hybrid concept has been successfully employed in the last decade for the development of Li+-doped ormolytes.9–24 The present work reviews the studies that have been carried out by our group for more than a decade to elucidate the structure, morphology, thermal properties and electrochemical behavior of a series of POE/ siloxane ormolytes doped with lithium salts. In an attempt to understand the ionic conductivity mechanism of the Li+-doped di-ureasils, we have used extensively vibrational spectroscopy to examine the cation/polymer, cation/anion and cation/crosslink interactions. The host matrix of the materials analysed belongs to the class of di-ureasils,25–27 a family of hybrids in which the organic and inorganic components are bonded through urea groups. The terminology adopted to represent these materials is the following:
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The non-doped di-ureasils are designated as d-U(Y), where d indicates di, U denotes the urea group and Y = 2000, 900 and 600 corresponds to the average molecular weight (in g mol−1) of the starting diamine precursor (corresponding to approximately 40.5, 15.5 and 8.5 (OCH2CH2) repeat units, respectively). The doped di-ureasils are described by the d-U(Y)nMXm notation, where n (salt composition) represents the number of ether oxygen atoms per Mm+ cation and X is the anion. For instance, the sample d-U(2000)20LiBF4 is a long chain di-ureasil incorporating an amount of lithium tetrafluoroborate such that the OPOE/Li+ ratio is equal to 20.
5.2
Ionic conductivity
The total ionic conductivity is an important characterization parameter which has been used as the criterion for possible application in devices. A PE intended for use in diverse electrochemical applications must have adequate ionic conductivity, together with negligible electronic conductivity if self-discharge on standing is to be avoided. A PE is considered to be a promising candidate for commercial application if its ionic conductivity is as high as 10−5 S cm−1 at room temperature.28–30 In general, the ionic conductivity of the electrolytes is measured as a function of salt composition and temperature. The objective of this characterization is to identify the electrolyte with the most favourable behavior for use as a component of the practical device. In general, salts with a polarizing cation and a large anion with a welldelocalized charge, and therefore also with low lattice energy, are the most suitable for use in PEs.31,32 In spite of the dangers associated with the anion, lithium perchlorate (LiClO4) is a salt that satisfies the conditions mentioned above. Lithium trifluoromethanesulfonate (or triflate) (LiCF3SO3) and LiBF4 have also been extensively employed in this context.3 Lithium bis(trifluoromethanesulfonyl)imide (LiTFSI) is particularly interesting as a guest species in solid PEs and also one of the best choices. In common with other salts that contain large polarizable anions, LiTFSI has low lattice energy and a low tendency to form ion-pairs, leading to enhanced ionic mobility. This salt also performs as a plasticizer in polyether electrolytes by creating free-volume. This is a significant advantage in polymer hosts that have an inherent tendency to crystallize. Figures 5.1–5.3 are included to demonstrate the effect of the choice of the di-ureasil network on the ionic conductivity. The Arrhenius plots show the variation of ionic conductivity with temperature of selected compositions of the d-U(2000)-, d-U(900)- and d-U(600)-based di-ureasil systems doped with the four lithium salts mentioned above.20,22,23,33
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5.1 Arrhenius conductivity plots of the d-U(2000)nLiX di-ureasil systems. Adapted from Silva et al.,20,21 Gomes Correia et al.23 and Barbosa et al.33
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5.2 Arrhenius conductivity plots of d-U(900)nLiX di-ureasil systems. Adapted from Barbosa et al.33–35
Typically the total ionic conductivity of the di-ureasils discussed here was measured by placing the sample between gold blocking electrodes, along the so-called electrode/di-ureasil/electrode assembly, which was secured in a suitable constant-volume support, to form a symmetrical cell.
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5.3 Arrhenius conductivity plots of d-U(600)nLiX di-ureasil systems. Adapted from Barbosa et al.33,35
Low-amplitude alternating potentials at frequencies between 65 kHz and 0.5 Hz were applied over a range of temperatures from 25 to 100 °C. It is clear from the analysis of Figs 5.1–5.3 that the materials characterized in this study show a non-linear variation of log conductivity with 1/T in the range between 25 and 100 °C. This behavior is characteristic of amorphous polymer electrolytes. The highest room temperature conductivity of the electrolyte system d-U(2000)nLiTFSI is 3.2 × 10−5 S cm−1, registered with the d-U(2000)35LiTFSI composition. This value is similar to that reported by Armand et al.36 for the POEnLiTFSI electrolyte system. The highest conductivity over the temperature range studied was found with the d-U(2000)10LiTFSI ormolyte, and was recorded as 1.2 × 10−3 S cm−1 at 95 °C (Fig. 5.1). The conductivity isotherms derived from conductivity measurements revealed the presence of a maximum conductivity for the electrolyte system located in the composition interval 8 ≤ n ≤ 10. Ionic conductivity in this hybrid electrolyte increases gradually with salt content for compositions with n ≥ 10. This parameter is generally expected to increase with the number of charge carriers, up to a limit found at about n = 8. Further increase in salt content beyond this concentration resulted in a decrease in total ionic conductivity. This observation can be explained by the formation of associated ionic species (e.g. contact ion pairs, triplets or higher multiplets or ultimately ion aggregates) and an increased tendency for ions to form bridging interactions between adjacent polymer chains a process
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designated as ‘ionic crosslinking’. Associated ionic species may be expected to show lower mobility than dissociated ions and the restriction of host matrix segment mobility also contributes to a reduction in the rate of ion transport and electrolyte conductivity. We will return to the aspect of ionic association in Section 5.5. For the hybrid system d-U(900)nLiTFSI, the most conducting composition is n = 15 which reaches 7.6 × 10−5 S cm−1 at 95 °C.34 At room temperature the sample with the highest conductivity is d-U(900)25LiTFSI (2.9 × 10−6 S cm−1). The conductivity isotherms of this electrolyte system confirmed that ionic transport in this hybrid electrolyte system is almost independent of the salt content over a wide range of composition. The broad conductivity maxima is located between n = 15 and 25. Further increase in salt content beyond about n = 15 led to a decrease in total ionic conductivity. This observation is consistent with a decrease in the flexibility of the host polymer chain segments, confirmed by an increase in the Tg of the electrolyte matrix. The Tg can be defined as the temperature at which the polymer chains in the amorphous phase begin segmental motion.3 According to the classical study of Berthier et al.5 it has been recognized that ionic mobility is closely correlated with the viscoelastic properties of the host polymer and so the relative magnitude of Tg in these polymer/salt complexes can be of considerable importance. As the mechanism of ion transport in amorphous PEs is dependent on the motion of polymer chain segments, the ionic conductivity is expected to decrease with an increase in Tg (Fig 5.4). The observed increase in Tg at lower values of n can be explained by the increase in number of bridging interactions which take place between adjacent polymer chains and guest salt species, generally designated as ‘ionic crosslinking’. Figure 5.2 compares the ionic conductivity of di-ureasil systems composed of the same host matrix. While at temperatures higher than 60 °C the d-U(900)nLiClO4 ormolytes are more conducting than comparable d-U(900)nLiTFSI materials, at lower temperatures the difference is negligible and LiTFSI is a better choice for safety reasons. The most conducting electrolyte of the hybrid d-U(900)nLiBF4 systems is the d-U(900)35LiBF4 composition (1.70 × 10−4 and 2.09 × 10−5 S cm−1 at about 95 °C and room temperature, respectively) (Fig. 5.2). These results are comparable with those reported by Chiodelli et al.37 and Armand et al.38 for the POEnLiBF4 system. The results show that for hybrid electrolyte networks conductivity increases with the salt content as the number of charged species available to support ionic conduction is increased and ionic transport is almost independent of the salt content over a wide range of composition of 40 ≥ n ≥ 200. In the case of the d-U(900)nLiBF4 system, isotherm curves reveal two conductivity maxima at n = 35 and at 20 ≥ n ≥ 10 at higher temperatures.
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5.4 Comparison of the variation of Tg and of the ionic conductivity of selected compositions of three di-ureasil systems doped with LiBF4 at 30 °C as a function of composition. Adapted from Barbosa et al.33
Figures 5.1–5.3 demonstrate that the di-ureasils doped with LiBF4 are the best conducting electrolytes. We may also infer from these data that the di-ureasils based on higher molecular weight POE chains are more conducting than ormolytes based on the di-ureasil structure d-U(600). These results were expected, since it is generally accepted that local segment motion plays an important role in the ionic migration in polymer electrolytes. The oxyethylene segments of d-U(600) are very short, causing restrictions of polymer segment mobility and a decrease in the transport of the guest ions. The higher molecular weight chains of the doped d-U(900) and d-U(2000) di-ureasils were found to support higher conductivity.20,33,35
5.3
Thermal properties
Thermal analysis techniques, in particular differential scanning calorimetry (DSC) and thermogravimetric analysis (TGA), are valuable tools to study the thermal behavior of PEs. DSC allows us to calculate the proportion of crystallinity, to detect the formation of new crystalline phases, free guest salt or uncomplexed polymer chains, to monitor the loss of solvent(s) (e.g. occluded water, alcohol), to determine the Tg value and to distinguish between endo- and exothermic events. TGA provides rich information
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n = 200 n = 80
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(b)
5.5 DSC curves of (a) d-U(2000)nLiTFSI (adapted from Barbosa et al.22) and (b) d-U(900)nLiTFSI (adapted from Barbosa et al.34) di-ureasil systems.
about the thermal degradation of the samples and their thermal stability domain. For the DSC measurements of di-ureasil ormolyte samples, sections have been usually removed from dry films and subjected to thermal analysis under a flowing inert atmosphere between 25 and 300 °C and at a heating rate of 5 °C min−1. Samples are then transferred to aluminium cans.
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Figure 5.5 reproduces the DSC curves obtained for the d-U(2000)nLiTFSI and d-U(900)nLiTFSI ormolytes, respectively.22,34 The absence of crystallinity in samples of both networks was expected from conductivity measurements. No endotherms associated with either the fusion of oxyethylene chain segments or the evaporation of water or ethanol are observed, leading us to conclude that the samples with 200 ≥ n ≥ 5 are completely amorphous. At high salt concentration (n = 0.5) an endothermic peak is observed at approximately 150 °C in the d-U(900)nLiTFSI system. This endothermic event is probably associated with the fusion of a salt/polymer complex, as reported in other di-ureasil systems based on other lithium salts.34 The onset of thermal decomposition was estimated by extrapolation from TGA curves (not shown). The lowest decomposition temperature observed with the d-U(900)nLiTFSI series was registered at n = 40 (358 °C), a finding that indicates that this di-ureasil system is more stable than the analogous d-U(900)nLiClO4 ormolyte. It is evident that the presence of the guest salt destabilizes the polymer matrix. However there is apparently no direct relationship between the onset of weight loss and the amount of salt present in the electrolyte. This situation is similar to that observed with electrolytes of other di-ureasil systems.20,35 Figure 5.4 illustrates the variation of Tg for selected compositions of the LiBF4-doped d-U(600)-, d-U(900)- and d-U(2000)-based electrolytes.33 The incorporation of guest salt leads to an increase in Tg, which means that the electrolyte sample becomes less flexible. Compositions with the highest salt content present poor mechanical properties, which in practical devices limits seriously the usefulness of the electrolyte in electrochemical applications. As expected from previous studies of similar POE/siloxane ormolytes,33 an increase in salt concentration causes a corresponding increase in the Tg of the di-ureasils (Fig. 5.4). Another aspect worth noting is that the increase in Tg with salt concentration is much more pronounced in the d-U(600)-based electrolytes than in the d-U(900)- or d-(2000)-based xerogels, which may be explained by the restrictions of segmental motion of the low molecular weight polyether chains present in the d-U(600) network. Figure 5.4 also reveals that the decrease in conductivity is consistent with an increasing Tg and significant levels of ionic conductivity exist only above this temperature.
5.4
Electrochemical stability
The application of the ormolytes in electrochemical applications depends on their stability window. To evaluate the electrochemical stability window of di-ureasil ormolyte compositions we have typically used a two-electrode cell configuration involving the use of a 25 µm-diameter gold microelectrode surface. Cell assembly was initiated by locating a freshly cleaned
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5.6 Voltammogram of d-U(2000)15LiTFSI di-ureasil electrolyte at a 25 µm diameter gold microelectrode vs. Li/Li+. Initial sweep direction is anodic and sweep rate is 100 mV s−1. Reproduced from Barbosa et al.22
lithium disk counter electrode on a stainless steel current collector. A thinfilm sample of ormolyte was centered over the counter electrode and the cell assembly completed by locating and supporting the microelectrode in the center of the electrolyte disk. Measurements were conducted at room temperature within a Faraday cage. The electrochemical stability range of the lithium-doped di-ureasils was determined by microelectrode cyclic voltammetry over the potential range between −1.5 and 6.5 V.20,22,33 In the anodic region, all ormolytes are stable up to 4.5 V vs. Li/Li+. Lithium deposition begins in the cathodic region at about −0.5 V vs. Li/Li+ (Fig. 5.6). These results show that the electrochemical stability of the di-ureasils is acceptable for application in practical devices.
5.5
Spectroscopic studies
Infrared and Raman spectroscopy are powerful tools in the elucidation of the ionic conductivity/ionic association relationship in PE systems. This sort of spectroscopic analysis usually involves the examination of diagnostic bands of the host polymer and of the anion. In the case of the organic/ inorganic hybrid matrices incorporating organic cross-links (e.g. diureasils), this study must also include the analysis of the bands characteristic of these functional groups (i.e. urea groups).
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The mid-infrared Fourier transform infrared (FT-IR) spectra (4000– 400 cm−1 range) should be acquired using a spectrometer placed inside a glove box with a dry argon atmosphere. If this is not possible, it is advisable to perform the measurement in vacuum, by averaging at least 150 scans at a high resolution (1–2 cm−1). The solid samples should be finely ground and mixed with dried spectroscopic grade potassium bromide. Prior to recording the spectra, it is highly recommended to keep the resulting pellets in an oven under vacuum at 90 °C for several days in order to reduce the levels of adsorbed water. The Fourier transform Raman (FT-Raman) spectra should be collected over the same wavenumber range, also at a high resolution (2–1 cm−1). An accumulation time for each spectrum of 4 hours is adequate. To evaluate complex band envelopes and to identify underlying component bands of the spectra, it is often necessary to perform classical curvefitting procedures, using for instance an iterative least-squares curve-fitting procedure such as that proposed in the Peakfit (San Rafael, CA) software. The best fit of the experimental data may be obtained by varying the frequency, bandwidth and intensity of the bands. In most cases it is appropriate to employ Gaussian band shapes.
5.5.1 Cation/polymer interactions Several vibrational modes of POE are sensitive to the interaction of the polymer backbone with cations and can be thus employed as diagnostic tools to monitor the changes undergone by the polyether chains upon addition of the guest salt. Two spectral regions have been widely used to probe the coordination of the cations to host polyethers: Skeleton COC stretching, νCOC, region. In this frequency interval (1200–1000 cm−1) the complexation of the cations by the heteroatoms of the polyether chains produces a shift of the strong νCOC band to lower wavenumbers. The magnitude of this shift depends on the strength of the interaction. • CH2 rocking, rCH2, region. Between 1000 and 800 cm−1 the bands ascribed to a mixture of CH2 rocking and CC stretching modes (1000– 900 cm−1) and the bands originated from CH2 rocking modes coupled with CO stretching modes (900–800 cm−1) appear.39–41 Their frequency strongly depends on the –O–C–C–O– angle and consequently on polymer conformation. •
A typical procedure adopted to monitor the coordination of cations to POE chains, based on the inspection of the skeleton νCOC modes in the FT-IR spectra, follows. The example given deals with a study previously carried out by our group on d-U(2000)nLiCF3SO3 di-ureasil ormolytes.17 In the
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5.7 FT-IR νCOC region of selected d-U(2000)nLiCF3SO3 di-ureasils (a) and results of the curve-fitting results for the most concentrated samples (b). Adapted from Nunes et al.17
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νCOC region the complexation of the Li+ ions by the oxygen atoms of the POE chains is reported to induce a shift of the prominent νCOC band to 1095–1089 cm−1.42,43 The room temperature FT-IR spectra of selected d-U(2000)nLiCF3SO3 samples in the νCOC region and the results of the curve-fitting performed with two representative samples are reproduced in Fig. 5.7. The FT-IR spectrum of d-U(2000) displays in this spectral region a prominent broad band around 1111 cm−1 and a shoulder at about 1148 cm−1 (Fig. 5.7), due to the νCOC vibration mode and to the coupled vibration of the νCO and rCH2 modes, respectively.39,40 The analysis of the spectra of the samples with 200 ≥ n ≥ 20 allowed us to conclude that the intensity and frequency of both features, characteristic of non-coordinated, disordered POE chains,39,40 remains essentially unchanged (Fig. 5.7). These results led us to conclude that in di-ureasil samples with n ≥ 20 the ether oxygen atoms of the polymer segments do not in principle bond to the Li+ ions. Their presence also made us suggest that the POE chains remain amorphous over the same composition range. Upon further addition of LiCF3SO3 to d-U(2000) (n = 10) the polymer features at 1148 and 1111 cm−1 persist, three new components emerge at 1162, 1137 and 1090 cm−1 and two new shoulders are seen at 1175 and 1080 cm−1 (Fig. 5.7). In the case of the d-U(2000)5LiCF3SO3 sample the νCOC band profile becomes considerably better resolved (Fig. 5.7). The 1175 and 1162 cm−1 bands – attributed to the asymmetric stretching vibration of the CF3 group – and the 1137, 1090 and 1080 cm−1 features coincide with those produced by the POE3LiCF3SO3 crystalline complex,44 indicating the formation of this compound in the diureasils with n = 10 and 5. The presence of the 1111 cm−1 band in the spectra of both samples supports that this crystalline compound coexists with noncoordinated and amorphous POE chains at n = 10 and 5. It must be emphasized that the determination of the composition that corresponds to the beginning of cation complexation by amorphous POE chains is often difficult. For instance, in the present case, although the first spectroscopic signs of Li+/POE bonding were evident in the νCOC region at n = 10, the salt concentration at which the Li+ ions start to effectively interact with the ether oxygen atoms of the POE chains may be lower. At n > 10 it is probable that the broad and strong νCOC envelope centered at 1111 cm−1 masks the characteristic band of Li+-coordinated amorphous POE chains (Fig. 5.7). To get more insight into the latter aspect it was considered of interest to examine the rCH2 region of the d-U(2000)nLiCF3SO3 ormolytes.45 The examination of the FT-IR spectra reproduced in Fig. 5.8 allowed us to immediately conclude that the rCH2 region of the Li+-based hybrids with 200 ≥ n ≥ 10 is very similar to that of d-U(2000). The main features of these spectra over the 980–900 cm−1 and 900–800 cm−1 ranges were a medium
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5.8 Room temperature FT-IR spectra of d-U(2000)nLiCF3SO3 di-ureasils in the rCH2 region. Reproduced from Nunes et al.45
intensity band at 950 cm−1 and a weak, broad and ill-defined band at 848 cm−1 (Fig. 5.8). Both features support that the conformation of the POE chains of d-U(2000) remains unaffected upon addition of increasing amounts of salt and therefore that the POE chains do not complex the Li+ ions at 200 ≥ n > 10. The spectra of the samples with n = 10, 5 and 1 differ markedly from those of the less concentrated materials. At n = 10 a feature at 969 cm−1, two medium intensity bands at 843 and 833 cm−1 and a weaker, broad event at 859 cm−1 emerge (Fig. 5.8). At higher salt concentration, the intensity of the bands increases significantly. At n = 1 the 937 and 843 cm−1 bands became the strongest events of the rCH2 region and a new band appears at 895 cm−1 (Fig. 5.8).
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Our interpretation of the spectral data obtained in the rCH2 region relied on early works of the infrared spectrum of high molecular weight poly(ethylene glycol)dimethyl ether (PEGDME).39–41 According to these studies, crystalline PEGDME exhibits a characteristic band at 844 cm−1,41 which undergoes a shift to 855 cm−1 upon melting,39,40 and a characteristic band at 937 cm−1.41 Consequently, we considered that the changes detected in the 1000–800 cm−1 interval of the FT-IR spectra of the d-U(2000)n LiCF3SO3 di-ureasils with 10 ≥ n ≥ 1 are a direct consequence of the alterations undergone by the POE chains to coordinate to the Li+ ions, since the POE segments of these three ormolytes adopt gauche conformations with a –O–C–C–O– torsional angle similar to that of the non-coordinated crystalline POE. In other words, in compounds with n > 10 the POE chains are amorphous, whereas in the samples with n = 10, 5 and 1 the chains adopt a helical-like structure. Comparison of the FT-IR spectrum of the POE3 LiCF3SO3 complex reported by Dissanayake and Frech44 with those of d-U(2000)5LiCF3SO3 and d-U(2000)1LiCF3SO3 revealed that the bands exhibited by the two hybrids at 969, 953, 937 and 833 cm−1 coincide with those produced by the crystalline POE3LiCF3SO3 complex, an indication of its formation in both salt-rich hybrid xerogels. A straightforward way of detecting the coordination of guest cations to the ether oxygen atoms of POE implies the analysis of the FT-Raman spectra in the region where the ‘oxygen breathing’ mode (νM–O) is expected (925–775 cm−1).46 In this spectral range the FT-Raman spectra of the d-U(2000)nLiCF3SO3 hybrids with n > 20 resemble that of d-U(2000) (Fig. 5.9). At n = 20 a shoulder emerges at 865 cm−1 and becomes a strong event in samples with n ≤ 7. Simultaneously the intensity of the 809 cm−1 band is reduced with the increase of salt concentration (Fig. 5.9). The 865 cm−1 band is attributed to a stretching vibration mode related with wrapping of the POE chains around the Li+ ion during complexation42,47,48 and therefore with the presence of gauche conformations of the O–C–C–O bonds,42 whereas the 809 cm−1 band indicates the presence of trans –O–C–C–O– conformations in the POE chains.49,50 The occurrence of the 865 cm−1 band in the FT-Raman spectra of the d-U(2000)nLiCF3SO3 with n ≤ 20 corroborates the claim that POE/Li+ interactions take place at high salt concentration. The conclusions derived from the νCOC17 and rCH245 regions allowed us to place the beginning of Li+/ether oxygen bonding near n = 10.
5.5.2 Cation/anion interactions In PEs several types of charge carriers may participate in the conduction mechanism: (a) ‘free’ or weakly coordinated ions with relatively high mobility; (b) cations interacting strongly with the host polymer and thus with low
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5.9 Room temperature FT-Raman spectra of d-U(2000)nLiCF3SO3 di-ureasils in the rCH2 region. Reproduced from Nunes et al.45.
mobility; (c) charged aggregates with moderate-low mobility. Papke et al.51 stated that in POE-based electrolytes containing monovalent ions an increase in the number of uncharged ion pairs is accompanied by a decrease in the ionic conductivity. To determine the extent of ionic association in PEs on the basis of vibrational spectroscopic studies, it is essential to introduce the cations as salts including appropriate anions. To be able to play the role of ion probes, the latter species must have vibration modes easily detected both in the infrared and Raman spectra, the attribution of which should be perfectly established and well documented in the literature.
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Three examples of these analyses will be given as follows. The first case mentioned here will be the spectroscopic study of the d-U(2000)-based diureasil system doped with LiCF3SO3.17,45 We will then refer a spectroscopic study of the d-U(2000)-based di-ureasil system doped with LiTFSI.52 Our third example is very recent and deals with d-U(2000)-based di-ureasil networks doped with LiBF4.53 Di-ureasils doped with LiCF3SO3 It is accepted that a ‘free’ CF3SO3− ion in a staggered configuration with a C3v point group symmetry has 18 normal modes with the symmetry representations 5A1+A2+6E. The A1 and E modes are infrared and Raman active. The A2 mode, associated with the internal torsion of the anion, is inactive both in infrared and Raman. Because all these modes are more or less sensitive to coordination effects they can be used to look into the anion local chemical surrounding.42,43,47,54–102 We will focus the following discussion on the non-degenerate symmetric stretching vibration mode of the SO3 group (νsSO3) of the CF3SO3− ion. Typically the νsSO3 band of a ‘free’ triflate ion appears at 1032 cm−1.103 Upon coordination to the Li+ ion, this feature shifts to higher wavenumbers.43,55,57,65,69,73,95–98,101,104,105 We will discuss first the νsSO3 region of the FT-IR spectra of the d-U(2000)nLiCF3SO3 compounds45 and then we will examine the information retrieved from the corresponding FT-Raman spectra.17 The FT-IR νsSO3 envelope and the results of the curve-fitting carried out are reproduced in Fig. 5.10(a) and (b), respectively. The presence of the 1031 cm−1 band in the FT-IR spectra of the doped samples indicates the presence of ‘free’ anions over the whole range of compositions studied. ‘Crosslink separated ion pairs’106 may also contribute to this feature. The shoulders at 1038 and 1025 cm−1 are attributed to weakly coordinated triflate ions located in two different anionic environments:106 (1) CF3SO3− species weakly bonded to Li+ ions, which simultaneously interact with the carbonyl oxygen atoms of the urea groups; (2) CF3SO3− ions hydrogenbonded to the urea N–H groups.106 The feature at 1043 cm−1 in samples with n ≤ 40 is associated with monodentate Li+CF3SO3− ions pairs or negatively charged triplets [Li(CF3SO3)2]− 57,104,105 The 1053 cm−1 component found at n = 5 and 1 is associated to a bidentate bridging aggregate (aggregate I)57 (the positively charged triplet [Li2(CF3SO3)]+).57,104,105 The event at 1063 cm−1 observed only in the case of the most concentrated di-ureasil is ascribed to a tridentate bridging aggregate (aggregate II)57 (the divalent positively charged multiplet [Li3(CF3SO3)]2+). The FT-Raman spectra of selected di-ureasils in the νsSO3 region and the results of the curve-fitting performed in the νsSO3 envelope are represented
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5.10 Room temperature FT-IR spectra (a) (adapted from Nunes et al.45), FT-Raman spectra (c) (adapted from Nunes et al.17) and corresponding curve-fitting results ((b) and (d), respectively) of selected d-U(2000)nLiCF3SO3 di-ureasils in the νsSO3 region. In order to examine exclusively the contribution of the νsSO3 mode, the FT-Raman spectrum of the pure polymer had to be first subtracted from those of the hybrids with n ≥ 10.
in Fig. 5.10(c) and (d), respectively. The detection of the 1032 cm−1 band in the FT-Raman spectra of all the doped di-ureasils represents more evidence of the existence of ‘free’ anions and very likely of the ‘crosslink separated ion pairs’106 over the whole range of salt concentration. This would explain the regular increase of the 1032 cm−1 feature as LiCF3SO3 concentration is increased and specially its marked intensity at n = 1. As noted above, the
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shoulders at 1039 and 1025 cm−1 are associated with weakly coordinated triflate ions located in two different sites. The component at 1045 cm−1 in samples with n ≤ 40 is associated with the formation of Li+CF3SO3− ion pairs or [Li(CF3SO3)2]− triplets.57,104,105 The component at 1052 cm−1 in the spectra of the di-ureasils with n = 5 and 1 is ascribed to the occurrence of the [Li2(CF3SO3)]+ triplet57,104,105 (aggregate I).52,57 This result is in perfect agreement with the identification of the POE3LiCF3SO3 crystalline complex65,69,73 at n = 5 and 1. In this compound the triflate ion vibrates essentially as the
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[Li2(CF3SO3)]+ entity because two oxygen atoms of the CF3SO3− ion bridge, in a monodentate fashion, two Li+ ions that lie adjacent to each other within the polymer helix.65,69,73 The event at 1063 cm−1 in the νsSO3 region of the most concentrated di-ureasil is associated with the formation of the [Li3(CF3SO3)]2+ multiplet (aggregate II).57 The anionic configurations detected in the FT-Raman νsSO3 region are in perfect agreement with those found in the FT-IR spectra. Comparison of the FT-IR45 and FT-Raman data17 of the d-U(2000)n LiCF3SO3 system allowed us to conclude that, although the number of components and their frequency coincide in samples with the same composition, their relative intensity differs considerably. For instance, while the 1031 cm−1 feature is the strongest event of the FT-Raman νsSO3 band of d-U(2000)5LiCF3SO3 (Fig. 5.10c), in the FT-IR spectrum of the same di-ureasil the most intense band is the 1043 cm−1 peak (Fig. 5.10a). We
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speculated that this mode is IR active, but weak in Raman scattering, meaning that the vibration associated with it may imply changes in the dipole moment but not in the polarizability of the species in question. Di-ureasils doped with LiTFSI The fundamental vibrations of the TFSI− ion are situated below 1400 cm−1.107 If the anion adopts C2 symmetry, its 39 internal vibrations can be classified into 20 A and 19 B modes, the former being polarized in Raman.107 In addition, for C2 symmetry, each of these modes will split into in-phase (A) and
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out-of-phase (B) components because of the coupling between the two SO2 groups.107 In our investigation of the d-U(2000)nLiTFSI materials we centered our attention on the region of the FT-Raman spectra characteristic of the symmetric deformation mode of the CF3 group (δsCF3) (also considered by some authors as a symmetric stretching mode of the SNS group (νsSNS)). We also examined the FT-IR spectra in the range of wavenumbers where the asymmetric stretching mode of the SNS group (or νaSNS), the asymmetric deformation mode of the CF3 group (δaCF3) and the δsCF3/νsSNS mode are expected. The Raman δsCF3 mode of the ‘free’ imide ion is found at 739 cm−1.107 This ion-pairing sensitive band shifts to higher wavenumbers upon coordination of the anion to a cation. We note that the assignment of this band has been a matter of debate. Bakker et al.108 reported that it has a marked νS–N character and therefore many authors prefer to associate it predominantly to the νsSNS.109,110 Henceforth this mode will be indicated as δsCF3/ νsSNS. Figure 5.11 reproduces the results of the curve-fitting carried out in the FT-Raman δsCF3/νsSNS band of LiTFSI-doped di-ureasils and the composition dependence of the area of the resolved components, respectively. At n ≥ 40 the band was decomposed into a single component at 739 cm−1 due to ‘free’ TFSI− ions,107 whereas at n = 20 it was resolved into two peaks at 743 and 739 cm−1. The former component is ascribed to the occurrence of contact ion pairs.107,108 Figure 5.11(b) shows that, while the concentrations of ‘free’ anions and coordinated anions remain essentially constant at 40 < n ≤ 8, at n = 5 the fraction of ‘free’ anions decreases significantly and the proportion of associated species increases concomitantly. The FT-IR spectra of the LiTFSI-doped di-ureasils with n ≥ 20 in the νaSNS, δaCF3 and δsCF3/νsSNS regions are shown in Fig. 5.12(a). The results of the curve-fitting performed in the νaSNS and δsCF3/νsSNS envelopes of selected samples are shown in Figs. 5.12(b) and (c), respectively. The FT-IR νaSNS region was decomposed into two components at 791 and 784 cm−1 for samples with 40 ≥ n ≥ 8 (Fig. 5.12b). The band at 784 cm−1 supports the presence of ‘free’ TFSI− ions,109 whereas that at 791 cm−1 is indicative of the occurrence of contact ion pairs. The FT-IR δsCF3/νsSNS band of the LiTFSIdoped di-ureasils with 40 ≥ n ≥ 8 was resolved into two components at 742 and 738 cm−1 (Fig. 5.12c). The 738 cm−1 band indicates the existence of ‘free’ anions,107–110 whereas the 742 cm−1 band is associated with the formation of contact ion pairs. The main conclusion derived from the study of the TFSI− characteristic modes of the d-U(2000)-based di-ureasils was that, although the Li+ ions interact extensively with the anions for compositions n ≤ 40, this process leads exclusively to the occurrence of contact ion pairs.
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5.11 Curve-fitting results of the FT-Raman δsCF3 region of selected d-U(2000)nLiTFSI di-ureasils (a) and salt concentration dependence of the integral intensity of the different spectral resolved components (b). The lines drawn are just guides for the eyes. Adapted from Nunes et al.17
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5.12 FT-IR spectra of selected d-U(2000)nLiTFSI of di-ureasils in the 800–700 cm−1 interval (a) and results of the curve-fitting of the νaSNS band (b) and δsCF3/νsSNS band (c) of representative samples. Adapted from Nunes et al.52
Di-ureasils doped with LiBF4 The non-bonded (‘free’) tetrahedral BF4− ion (point group Td) has nine vibration normal modes: the non-degenerate mode ν1(A1); the doubly degenerate mode ν2(E); the triply degenerate modes ν3(T2) and ν4(T2). All the modes are Raman active, but only ν3(T2) and ν4(T2) are infrared active. The lowering of the local symmetry around BF4− resulting from coordination of this anion to a metal cation has the following consequences: (i) band splitting of degenerate vibrations; (ii) frequency shifts of non-degenerate vibrations; (iii) activation of the infrared-forbidden ν1 and ν2 modes. The magnitude of the band splittings and of the frequency shifts depends on the strength of ion association. The BF4− ion may adopt three possible coordinating configurations: monodentate, tridentate or bidentate. •
Monodentate (or tridentate) configuration. The association of one (or three) of the BF4− fluorine atoms with a cation lowers the anion Td
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symmetry of the anion to C3v. Both configurations are manifested through the splitting of the threefold degenerate ν3 and ν4 vibrations into two components (A1 and E) both in the infrared and Raman spectra. • Bidentate configuration. The association of a cation with two of the BF4− fluorine atoms reduces the anion symmetry to C2v. This configuration leads to the splitting of the twofold degenerate ν2 vibration into two components (A1 and A2) in the Raman spectrum and to the splitting of the threefold degenerate ν3 and ν4 vibrations into three components (A1, B1 and B2) both in the infrared and Raman spectra. Theoretical quantum chemical calculations suggested that the bidentate structure is the most favourable for ion pairs of BF4− with Li+.111–113
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In the Raman spectrum of molten/aqueous solution of sodium tetrafluoroborate (NaBF4) the ν1(A1), ν2(E), ν3(T2) and ν4(T2) modes appear at 777/773 cm−1, 360/357 cm−1, 1070/1080 cm−1 and 533/528 cm−1, respectively.114 The appearance of the IR inactive ν1 mode at 770 cm−1 in the FT-IR spectra of the LiBF4-doped di-ureasils shown in Fig. 5.13 revealed the formation of ion pairs over the whole range of salt composition studied. The three features seen at 574, 534 and 522 cm−1 in the same spectra were tentatively attributed to the lifting of the degeneracy of the ν4BF4 mode resulting from the formation of ion pairs in a presumably bidentate coordination. This spectroscopic analysis enabled us to conclude that in all the d-U(2000)nLiBF4 samples examined there was a considerable tendency for ionic association.
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5.13 FT-IR spectra of selected d-U(2000)nLiBF4 di-ureasils in the ν1BF4− and ν4BF4− regions. Reproduced from Barbosa et al.35
5.5.3 Cation/crosslink interactions In di-ureasil-based ormolytes doped with a lithium salt three coordinating sites are available for the Li+ ions: the urea carbonyl oxygen atoms (–C=O) and the polymer ether oxygen atoms (–COC–) of the host di-ureasil matrix and the coordinating atoms of the guest anion (e.g. the triflate oxygen atoms (CF3SO3−) in the case of LiCF3SO3, the perchlorate oxygen atoms (ClO4−) in the case of LiClO4, etc.). To illustrate the spectroscopic analysis usually adopted to study Li+/urea interactions in di-ureasil systems, we will refer in this section to the analysis carried out with the d-U(2000)-based xerogels doped with LiCF3SO3.45
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Data retrieved from other spectral regions had demonstrated that in d-U(2000)nLiCF3SO3 samples with n > 20: (1) the POE chains do not complex the Li+ ions and (2) the concentration of coordinated anionic species is quite low. Therefore it seemed obvious to assume that at n > 20 the coordination of the Li+ ions in the d-U(2000)-based framework is entirely ensured by the C=O groups. At n > 20 two sorts of interactions could be then envisaged for the Li+ ions after being incorporated into the di-ureasil framework: they could either bond to ‘free’ C=O groups or coordinate to hydrogen-bonded C=O groups. The latter situation is more complex since the Li+ ions need to disrupt the hydrogen-bonded array formed throughout the d-U(2000) framework via urea groups. The ability of urea groups to participate in highly specific hydrogen bonding interactions is well known. In bis-ureido compounds, these interactions are responsible for unique gelling properties.115,116 This behavior relies essentially on geometrical considerations, since the C=O group of each urea group may form two hydrogen bonds with the N–H moieties of a neighbor urea group, leading to the formation planar bifurcated hydrogen bonds.117 In non-doped di-ureasils disordered POE/urea hydrogen-bonded associations and highly ordered self-associated urea/urea hydrogen-bonded structures were identified.118 To examine not only the Li+/urea interactions but also the strength and extent of hydrogen bonding in the d-U(2000)nLiCF3SO3 xerogels, we decided to inspect the ‘amide I’ and ‘amide II’ bands. The FT-IR spectra of selected samples in both spectral regions are reproduced in Fig. 5.14(a). ‘Amide I’ region (1800–1600 cm−1). The ‘amide I’ region of the di-ureasils corresponds to the amide I119 region of polyamides.120 The amide I mode (or simply carbonyl stretching mode) is a very complex vibration that receives a major contribution from the C=O stretching vibration.120 As it is sensitive to the specific type and magnitude of hydrogen bonding, the amide I band consists of several distinct components which correspond to different C=O environments, usually designated as associations, aggregates or structures. Since the absorption coefficients of C=O groups belonging to different aggregates may be different, only the changes undergone by each mode provide information of concentration variations of each aggregate.120,121 Figure 5.14(a) shows that the addition of Li+ cations to d-U(2000) perturbs deeply the ‘amide I’ region, an indication that the Li+ ions coordinate to the carbonyl oxygen atoms of the urea crosslinks over the whole range of salt concentration examined. The intensity maximum of the ‘amide I’ envelope, situated around 1721 cm−1 in samples with n ≥ 40, is displaced to 1662 cm−1 for n < 40. This result suggests that the hydrogen-bonded
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associations formed at high salt concentration are more ordered (i.e. stronger) than those found in the more dilute xerogels. To get additional insight into the modifications that take place in the ‘amide I’ region of the d-U(2000)nLiCF3SO3 materials at increasing salt concentration, we performed curve-fitting in the 1800–1600 cm−1 envelope. The results obtained are illustrated in Fig. 5.14(b). The composition dependence of the area of the resolved components is shown in Fig. 5.14(c).
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The ‘amide I’ band of d-U(2000) was decomposed into five individual components situated at 1751, 1720, 1686, 1662 and 1640 cm−1 (Fig. 5.14b). The band at 1751 cm−1 is associated with the absorption of urea groups in which the N–H or C=O groups are free from any interactions.118 The bands at 1720, 1686 and 1662 cm−1 are ascribed to hydrogen-bonded C=O groups of disordered hydrogen-bonded POE/urea associations of increasing strength.118 The 1640 cm−1 feature is assigned to the absorption of C=O groups included in more ordered hydrogen-bonded urea/urea associations.118
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Globally, the increase of LiCF3SO3 content has three major consequences (Figs. 5.14b and c): The 1752, 1720 and 1642 cm−1 features, ascribed to free C=O groups, disordered POE/urea aggregates and highly ordered urea/urea associations, respectively, suffer a considerable loss of intensity and vanish at n = 5. • The intensity of the 1662 cm−1 band, due to ordered POE/urea structures, is significantly enhanced. • Two new events emerge at 1710 and 1632 cm−1 in the case of the salt-rich materials with n = 5 and 1. •
The most relevant effects, observed at n = 5 (i.e. the growth of new components at 1710 and 1632 cm−1 and the concomitant disappearance of the 1752, 1720 and 1642 cm−1 events), led us to conclude, not only that the saturation of the urea crosslinks is attained at this composition (i.e. no C=O groups remain free), but also that new disordered POE/urea aggregates (1710 cm−1) and highly ordered hydrogen-bonded aggregates (1631 cm−1) are formed at the expense of the disruption of the more disordered POE/urea structures of d-U(2000) (1720 cm−1) and of the highly ordered urea/urea aggregates that give rise to the 1642 cm−1 component. At n = 1 a marked increase of the area of the 1631 cm−1 feature and a sharp reduction of the proportion of the POE/urea aggregates (1662 cm−1) occur. The latter effect must be correlated with the fact that at this salt content the POE segments are
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massively required to complex the Li+ ions to form the POE3LiCF3SO3 compound and can thus no longer interact with the N–H groups of the urea linkages through hydrogen-bonding to form such hydrogen-bonded aggregates. ‘Amide II’ region The ‘amide II’ mode of the di-ureasils corresponds to the amide II119 feature of polyamides.120 The amide II (or N–H in-plane bending) band is a mixed mode containing a major contribution from the N–H in-plane bending vibration.120 It is sensitive to both chain conformation and intermolecular hydrogen bonding and provides rich information about the distribution of hydrogen bond strengths.120 In the spectrum of d-U(2000) the ‘amide II’ mode appears as a medium intensity, broad band centered at about 1560 cm−1 (Fig. 5.14a),118 which was decomposed into three components at 1571, 1555 and 1529 cm−1 (Fig. 5.14b). This suggests that in the host hybrid hydrogen-bonded aggregates with at least three distinct degrees of order exist. The inclusion of LiCF3SO3 affects the ‘amide II’ envelope. At n = 20 the band profile is significantly modified, because of the loss of the 1529 cm−1component and the growth of a new component at 1586 cm−1, which shifts the intensity maximum of the band envelope to 1571 cm−1 (Fig. 5.14b). This finding indicates that, in agreement with the conclusions retrieved from the analysis of the ‘amide I’ region, in the samples with n ≤ 20 the hydrogen bonds are stronger than in the less concentrated ones.118
5.6
Electrochromic displays
Electrochromic (EC) technology and EC devices are drawing attention for their high potential with respect to solar control and display applications. EC materials have received a great deal of attention due to the important applications stemming from such systems, which include smart windows, displays, anti-glare mirrors, eye-glasses or solar-attenuated windows.122,123 Electrochromism may be defined as the phenomenon that gives rise to a reversible color change by the electrochemical insertion and extraction of electrons and ions into inorganic materials, for example, tungsten oxide (WO3). The difference in the transmittance between the colored (written) state and the fully bleached (unwritten) state indicates the light modulation capability of ECDs. A PE component in an optical device must meet various prerequisites. One of these is that the film must exhibit high transmission. If this condition is not fulfilled, the electrolyte will reduce the color contrast of the optical device.124
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The central part of an ECD is the electrolyte or ion conductor. It can be made of a thin (inorganic) film, a polymeric layer or a liquid capable of ion transport. PEs are of great and increasing interest for ECDs and allow convenient design incorporating resilient and adhesive layers. In contrast with the liquid electrolyte, there are no obvious problems with leakage and fluid-mechanics-induced distortions.125 In this section we describe the assembly of a prototype solid-state ECD based on a four-layer sandwich structure incorporating materials based on di-ureasils doped with controlled quantities of lithium salts.22,33,35,126 The device assembly was carried out by direct application of a small volume of the di-ureasils to the surface of a glass plate onto which an indium zinc oxide (IZO)/WO3 coating had been previously deposited. The gel was dried in a vacuum oven for a period of seven days. The thickness used for each layer were 170 nm for IZO and 400 nm for WO3. The entire assembly procedure described was carried out under atmospheric condition. A second IZO-coated glass plate was placed on top of the dry electrolyte layer and the two plates were pressed together to spread the electrolyte in a thin film between the EC surfaces. The entire assembly procedure described was carried out under atmospheric conditions. Figure 5.15 reports the optical transmittance in the 300–900 nm wavelength range for the ECDs based on d-U(900)5LiTFSI and d-U(600)25LiBF4 di-ureasils. These images show the compositions with good optical performance and stability. The average transmittance in the visible region of the spectrum was above 69 and 74% for devices based on d-U(900)5LiTFSI and d-U(600)25LiBF4 di-ureasils, respectively. After coloration, while the device based on d-U(900)5LiTFSI presented a good color contrast (above 24%) and an optical density above 0.49, that based on d-U(600)25LiBF4 exhibited an average transmittance of 48% and an optical density above 0.17, providing a good performance in the coloring/bleaching process (Fig. 5.16). These results are very satisfactory, especially when compared to those obtained with the incorporation of d-U(900)nLiClO4 electrolytes.35 The reasons for the variation of optical density with electrolyte composition are not yet clear and require further study. These findings are consistent with mechanical interfacial effects or the supply of ions due to higher ionic conductivity, both aspects of electrolyte performance that vary significantly with composition. The solid-state ECD tested exhibit good electrochemical stability, showed memory effect and may therefore be of interest for application in ‘smart windows’. The evaluation of devices based on di-ureasils confirmed that hybrid electrolytes provide adequate optical density, good stability and enhanced optical memory. Therefore PEs based on di-ureasil composites incorporating lithium guest salts represent an attractive alternative to liquid electrolytes in ECDs.126
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5.15 Optical transmittance as a function of wavelength for the electrochromic device in bleached and colored state using (a) d-U(900)5LiTFSI and (b) d-U(600)25LiBF4 di-ureasils. Adapted from Rodrigues et al.126
5.7
Conclusion
The present chapter, devoted to a series of attractive lithium ion conductors based on di-urea crosslinked POE/siloxane host networks (di-ureasils) prepared by the sol–gel process, demonstrates that the combination of the hybrid strategy with the host–guest concept derived from inorganic chemistry is extremely interesting. We have shown that the ormolytes discussed here exhibit good thermal, mechanical and electrochemical properties. In particular, we have shown that these materials can be employed as the ion conducting layer of practical
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5.16 ECD in bleached and colored states for selected di-ureasils: (a) d-U(900)5LiTFSI; (b) d-U(600)25LiBF4. Adapted from Rodrigues et al.126
EDCs. The prototype solid state ECDs assembled present encouraging optical contrast and appropriate response times. The advantages of using solid PEs instead of conventional liquid-based electrolytes in ECDs are related to the capacity of preventing leaking, eliminating the need for sealing, since no solvent evaporation takes place. Moreover, ormolytes are flexible materials and thus able to provide good contact between the electrolytic component and the active EC layer. In addition, film deposition of large surface areas is possible, while ensuring that the device is mechanically robust. Organic–inorganic hybrid materials are being increasingly explored and more developments are expected within the next years. It is worth noting that many of these hybrids satisfy socio-economical demands in energy and the environment.
5.8
References
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72. york s, kellam iii e c, allcock h r and frech r, ‘A vibrational spectroscopic study of lithium triflate in polyphosphazenes with linear oligoethyleneoxy sidechains of different lengths’, Electrochim Acta, 2001, 46, 1553–1557 73. rhodes c p and frech r, ‘Local structures in crystalline and amorphous phases of diglyme–LiCF3SO3 and poly(ethylene oxide)–LiCF3SO3 systems: implications for the mechanism of ionic transport’, Macromolecules, 2001, 34, 2660–2666 74. papke b l, ratner m a and shriver d f, ‘Vibrational spectroscopy and structure of polymer electrolytes, poly(ethylene oxide) complexes of alkali metal salts’, J Phys Chem Solids, 1981, 42, 493–500 75. macfarlane d r, meakin p, bishop a, mcnaughton d, rosalie j m and forsyth m, ‘FTIR study of ion-pairing effects in plasticized polymer electrolytes’, Electrochim Acta, 1995, 40(13–14), 2333–2337 76. bishop a, macfarlane d r, mcnaughton d and forsyth m, ‘FT-IR investigation of ion association in plasticized solid polymer electrolytes’, J Phys Chem, 1996, 100, 2237–2243 77. bishop a, macfarlane d r and mcnaughton d, ‘Triflate ion association in plasticized polymer electrolytes’, Solid State Ionics, 1996, 85, 129–135 78. ferry a, jacobsson p and torell l m, ‘The molar conductivity behavior in polymer electrolytes at low salt concentrations; a Raman study of poly(propylene glycol) complexed with LiCF3SO3’, Electrochim Acta, 1995, 40(13–14), 2369–2373 79. ferry a, jacobsson p and stevens j r, ‘Studies of ionic interactions in poly(propylene glycol)4000 complexed with triflate salts’, J Phys Chem, 1996, 100, 12574–12582 80. ferry a, orädd g and jacobsson p, ‘A pfg-NMR investigation of the diffusion coefficients of anionic species in the system PPG4000–LiCF3SO3’, Macromolecules, 1997, 30, 7329–7331 81. ferry a and tian m, ‘Influence of hydroxyl terminal groups on the ionic speciation and ionic conductivity in complexes of poly(propylene glycol)(4000) and LiCF3SO3 salt’, Macromolecules, 1997, 30, 1214–1215 82. ferry a, ‘Effects of dynamic spatial disorder on ionic transport properties in polymer electrolytes based on poly(propylene glycol)(4000)’, J Chem Phys, 1997, 107(21), 9168–9175 83. ferry a, ‘Ionic interactions and transport properties in methyl terminated poly(propylene glycol)(4000) complexed with LiCF3SO3’, J Phys Chem B, 1997, 101, 150–157 84. ferry a, doeff m m and dejonghe l c, ‘Transport property and Raman spectroscopic studies of the polymer electrolyte system P(EO)n-NaTFSI’, J Electrochem Soc, 1998, 145, 1586–1592 85. ferry a, orädd g and jacobsson p, ‘A Raman, ac impedance and pfg-NMR investigation of poly(ethylene oxide) dimethyl ether (400) complexed with LiCF3SO3’, Electrochim Acta, 1998, 43(10–11), 1471–1476 86. ferry a, edman l, forsyth m, macfarlane d r and sun j, ‘Connectivity, ionic interactions, and migration in a fast-ion-conducting polymer-in-salt electrolyte based on poly(acrylonitrile) and LiCF3SO3’, J Appl Phys, 1999, 86(4), 2346–2348 87. edman l, ferry a and jacobsson p, ‘Effect of C60 as a filler on the morphology of polymer–salt complexes based on poly(ethylene oxide) and LiCF3SO3’, Macromolecules, 1999, 32, 4130–4133
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88. stevens j r and shantz s, ‘Variation of the glass transition temperature of poly(propylene glycol)4000 complexed with lithium and sodium salts’, Polym Commun, 1988, 29, 330–331 89. shantz s, sandahl j, börjesson l, torell l m and stevens j r, ‘Ion pairing in polymer electrolytes; a comparative Raman study of NaCF3SO3 complexed in poly(propylene-glycol) and dissolved in acetonitrile’, Solid State Ionics, 1988, 28–30, 1047–1053 90. kakihana m, schantz s and torell l m, ‘Raman spectroscopic study of ion–ion interaction and its temperature dependence in a poly(propylene-oxide)-based NaCF3SO3–polymer electrolyte’, J Chem Phys, 1990, 92(10), 6271–6277 91. kakihana m, schantz s, torell l m and stevens j r, ‘Dissociated ions and ion–ion interactions in poly(ethylene oxide) based NaCF3SO3 complexes’, Solid State Ionics, 1990, 40–41, 641–644 92. shantz s, torell l m and stevens j r, ‘Ion pairing effects in poly(propylene glycol)–salt complexes as a function of molecular weight and temperature: a raman scattering study using NaCF3SO3 and LiClO4’, J Chem Phys, 1991, 94(10), 6862–6867 93. yang x q, lee h s, hanson l, mcbreen j and okamoto y, ‘Development of a new plasticizer for poly(ethylene oxide)-based polymer electrolyte and the investigation of their ion-pair dissociation effect’, J Power Sources, 1995, 54, 198–204 94. johnston s f, ward i m, cruickshank j and davies g r, ‘Spectroscopic studies of triflate ion association in polymer gel electrolytes and their constituents’, Solid State Ionics, 1996, 90, 39–48 95. torell l m, jacobsson p and petersen g, ‘A Raman study of ion solvation and association in polymer electrolytes’, Polym Adv Tech, 1992, 4, 152–163 96. petersen g, torell l m, panero s, scrosati b, silva c j and smith m j, ‘Ionic interactions in MCF3SO3-polyether complexes containing mono-, di- and trivalent cations’, Solid State Ionics, 1993, 60, 55–60 97. brodin a, mattson b, nilsson k, torell l m and hamara j, ‘Ionic configurations in PEO based electrolytes endcapped by CH3-, OH- and SO3-groups’, Solid State Ionics, 1996, 85, 111–120 98. petersen g, brodin a, torell l m and smith m j, ‘Light scattering and luminescence studies of M(CF3SO3)x–polyether complexes containing trivalent cations’, Solid State Ionics, 1994, 72, 165–171 99. brodin a, mattsson b and torell l m, ‘A site selective luminescence excitation and Raman scattering study of Eu(CF3SO3)3-PPO; ion association and ionic exchange rate’, J Chem Phys, 1994, 101(6), 4621–4627 100. bernson a and lindgren j, ‘Ion aggregation and morphology for poly (ethylene oxide)-based polymer electrolytes containing rare earth metal salts’, Solid State Ionics, 1993, 60, 31–36 101. bernson a and lindgren j, ‘Free ions and ion pairing/clustering in the system LiCF3SO3–PPOn’, Solid State Ionics, 1993, 60, 37–41 102. gejji s p, hermansson k, tegenfeldt j and lindgren j, ‘Geometry and vibrational frequencies of the lithium triflate ion pair: an ab initio study’, J Phys Chem, 1993, 97, 11402–11407 103. wendsjö å, lindgren j, thomas j o and farrington g c, ‘The effect of temperature and concentration on the local environment in the system M(CF3SO3)2PEOn for M = Ni, Zn and Pb’, Solid State Ionics, 1992, 53–56, 1077
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104. stevens j r and jacobsson p, ‘A comparison of acetone and poly(propy1ene glycol) as solvents for lithium triflate and lithium perchlorate’, Can J Chem, 1991, 69, 1980–1984 105. jacobsson p, albisson i, mellander b-e and stevens j r, ‘Ion association effects and phase separation in poly(propylene oxide) modified poly(dimethylsiloxane) complexed with triflate salts’, Polymer, 1992, 33(13), 2778–2783 106. de zea bermudez v, ostrovskii d, lavoryk s, gonçalves m c and carlos l d, ‘Urethane cross-linked poly(oxyethylene)/siliceous nanohybrids doped with Eu3+ ions. Part II: Ion association’, Phys Chem Chem Phys, 2004, 6(3), 649–658 107. rey i, johansson p, lindgren j, lassègues j c, grondin j and servant l, ‘Spectroscopic and theoretical study of (CF3SO2)2N− (TFSI−) and (CF3SO2)2NH (HTFSI)’, J Phys Chem A, 1998, 102(19), 3249–3258 108. bakker a, gejji s, lindgren j, hermansson k and probst m m, ‘Contact ion pair formation and ether oxygen coordination in the polymer electrolytes M[N(CF3SO2)2]2PEOn for M = Mg, Ca, Sr and Ba’, Polymer, 1995, 36, 4371–4378 109. ferry a, furlani m, franke a, jacobsson p and mellander b-e, ‘Spectroscopic studies of luminescent and ionically conducting Eu[N(CF3SO2)2]3-PPG4000 complexes’, J Chem Phys, 1998, 109(7), 2921–2928 110. furlani m, ferry a, franke a, jacobsson p and mellander be, ‘Time resolved luminescence and vibrational spectroscopic studies on complexes of poly(ethylene oxide) oligomers and EuTFSI3 salt’, Solid State Ionics, 1998, 113–115, 129–138 111. francisco j s and williams i h, ‘Structural and spectral consequences of ion pairing. 4. Theoretical study of alkali metal tetrafluoroborates (M+BF4−; M = lithium, sodium, potassium, and rubidium)’, J Phys Chem, 1990, 94, 8522–8529 112. suresh c h, gadre s r and gejji s p, ‘Theoretical studies on the structure of M+ BF4− ion pairs M = Li+, NH4+: the role of electrostatics and electron correlation’, Theor Chem Acc, 1998, 99, 151–157 113. dhumal n r. and gejji s p, ‘Theoretical studies on blue versus red shifts in diglyme–M+–X– (M = Li, Na, K and X = BF4, ClO4, SCN)’, J Molec Struct (THEOCHEM), 2006, 758, 233–240 114. quist a s, bates j b and boyd g h, ‘Raman spectra of molten NaBF4 to 606 °C and 8% NaF-92% to 503 °C’, J Chem Phys, 1971, 54, 4896–4901 115. de loos m, van esch j, stokroos i, kellogg r m and feringa b l, ‘Remarkable stabilization of self-assembled organogels by polymerization’, J Am Chem Soc, 1997, 119, 12675–12676 116. van esch j, kellogg r m and feringa b l, ‘Di-urea compounds as gelators for organic solvents’, Tetrahedron Lett, 1997, 38(2), 281–284 117. born l and hespe h, ‘On the physical crosslinking of amine-extended polyurethane urea elastomers: a crystallographic analysis of bis-urea from diphenyl methane-4-isocyanate and 1,4-butane diamine’, Colloid Polym Sci, 1985, 263, 335–341 118. de zea bermudez v, carlos l d and alcácer l, ‘Sol–gel derived urea crosslinked organically modified silicates. 1. Mid-infrared spectra’, Chem Mater, 1999, 11, 569–580 119. miyazawa t, shimanouchi t and mizushima s-i, ‘Characteristic infrared bands of monosubstituted amides’, J Chem Phys, 1956, 24(2), 408–418
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120. skrovanek d j, howe s e, painter p c and coleman m m, ‘Hydrogen bonding in polymers: infrared temperature studies of an amorphous polyamide’, Macromolecules, 1985, 18, 1676–1683 121. coleman m m, lee k h, skrovanek d j, and painter p c, ‘Hydrogen bonding in polymers. 4. Infrared temperature studies of a simple polyurethane’, Macromolecules, 1986, 19, 2149–2157 122. wang x j, lau w m and wong k y, ‘Display device with dual emissive and reflective modes’, Appl Phys Lett, 2005, 87, 113502-1–113502-3 123. mortimer r g, ‘Electrochromic materials’, Chem Soc Rev, 1997, 26, 147–156 124. cui m-h, guo j-s, xie h-q, wu z-h and qiu s-c, ‘All-solid-state complementary electrochromic windows based on the oxymethylene-linked polyoxyethylene complexed with LiClO4’, J Appl Polym Sci, 1997, 65, 1739–1744 125. granqvist c g, azens a, hjelm a, kullman l, niklasson g a, rönnow d, strømme mattsson m, veszelei m and vaivars g, ‘Recent advances in electrochromics for smart windows applications’, Solar Energy, 1998, 63, 199–216 126. rodrigues l c, barbosa p c, silva m m, smith m j, gonçalves a and fortunato e, ‘Application of hybrid materials in solid-state electrochromic devices’, Opt Mater, 2009, 31, 1467–1471
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6 Hybrid inorganic–organic polymer electrolytes V. D I N O T O, E. N E G R O and S. L AV I N A, University of Padova, Italy, and M. V I T TA D E L L O, City University of New York, USA
Abstract: Polymer electrolytes (PEs) are macromolecular systems capable of transporting charged species such as ions or protons. The main application of PEs is in energy conversion and storage devices such as batteries and fuel cells. The chapter overviews the synthesis, structure, physical and electrical properties of three classes of hybrid inorganic– organic PEs: three-dimensional hybrid inorganic–organic networks as polymer electrolytes (3D-HION-APE), zeolitic inorganic–organic polymer electrolytes (Z-IOPEs) and hybrid gel electrolytes (HGEs). The basic structure of the materials consists of organic macromolecules bridging inorganic clusters or species. The chapter also includes an overview of the methods used in the characterization of the structure and of the electrical conductivity of PEs, with a particular reference to the jump relaxation model. Key words: inorganic–organic polymer electrolytes, Z-IOPE, electrical spectroscopy, vibrational spectroscopy, conductivity spectra.
6.1
Introduction
One of the most active research areas in solid state electrochemistry concerns the development of ion-conducting materials for application in the conversion and storage of energy (e.g. high energy density batteries and fuel cells). Polymer electrolytes (PEs) are a class of materials which are particularly promising in this sense, as witnessed by the massive research efforts spent in the last 20 years to obtain systems characterized by a good conductivity and a high chemical, thermal and electrochemical stability. This chapter overviews hybrid inorganic–organic PEs, which are characterized by inorganic species included in the chemical composition of the materials and/or in their structure at a molecular or nanometric level. Three families of hybrid inorganic–organic PEs are reviewed: (a) three-dimensional hybrid inorganic–organic networks as polymer electrolytes (3D-HION-APEs); (b) zeolitic inorganic–organic polymer electrolytes (Z-IOPEs); and (c) hybrid gel electrolytes (HGEs). It is shown that these materials are able to transport either monovalent species (e.g. lithium ions or protons) or bivalent species (e.g. magnesium ions), showing a remarkable conductivity (ca. 10−5 S cm−1 or more at room temperature). In addition, nanocomposite 219 © Woodhead Publishing Limited, 2010
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inorganic–organic proton-conducting PEs based on perfluorinated ionomers and various ceramic oxoclusters are briefly overviewed. The chapter includes a discussion about the methods used in the structural characterization and in the study of the conductivity of these families of hybrid inorganic–organic PEs. In conclusion, the jump relaxation model is discussed as a theoretical framework for the interpretation of the electrical response of PEs in terms of the real component of their conductivity spectra.
6.2
Fundamentals of polymer electrolytes
PEs are crucial materials for the development of modern electric devices such as high energy density batteries, sensors, fuel cells and electrochromic displays (Scrosati, 1993, p182; Gray, 1997; Wright, 1998; Dias et al., 2000). After more than 20 years of research in the field of PEs, the distinct desirability of an entirely solid state energy storage system is still driving the quest to enhance the ionic conductivity of ‘dry’ solid PEs (higher than 10−3 S cm−1 at room temperature), as described by Scrosati (1993, p182), Gray (1997) and Armand et al. (2002). This is still true even after the recent introduction of the lithium–polymer technology for the reversible storage of energy. The latter systems are based on polymer gels consisting of liquid electrolytes immobilized in various polymer matrices, resulting in a conductivity of 10−2–10−3 S cm−1 at room temperature (Feuillade and Perche, 1975; Abraham and Alamgir, 1990; Croce et al., 1993; Stallworth et al., 1995). Unfortunately, these latter systems have the following drawbacks: (a) leakage of liquid upon squeezing and (b) loss of specific power and energy due to the inert host polymer. Up to now, however, the early expectation to obtain electrolytic complexes with conductivities comparable to those of super-ionic conductors has not been met, as discussed by Gadjourova et al. (2001). The research in the field of PEs is now facing a major dualism regarding the best way to improve ionic conductivity in conventional systems. On one hand, many highly amorphous electrolytic complexes have been synthesized with the goal of reducing the crystalline order as much as possible. This approach included the use of inorganic–organic polymer hosts with a low glass transition temperature (Tg), as discussed by Bouridah et al. (1985) and Blonsky et al. (1986), and the addition of plasticizers to polyether electrolytes (Huq et al., 1992). On the other hand, more recently it was suggested that the design of preferentially ordered materials could provide faster ion-conducting pathways. This has been achieved in liquid crystalline PEs (Wright et al., 1998) and in stretched PEs above the Tg (Chung et al., 1999). Finally, proton-conducting PEs play a crucial role in the development of advanced energy conversion devices such as proton exchange membrane fuel cells (PEMFCs). These systems are particularly attractive
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as sources of power in portable applications and light-duty vehicles because of their high energy density (Cleghorn et al., 1997; Appleby, 1999; Larminie and Dicks, 2000; Ogden, 2003; Theisen, 2003). Perfluorinated polymer electrolytes such as Nafion, Aciplex, Flemion and Dow membranes are some of the most promising electrolyte membranes for PEMFCs (Mauritz and Moore, 2004). Nevertheless, the major drawbacks to their large-scale commercial use involve the high costs, the large crossover of reagents and a low proton conductivity at a low hydration degree and at temperatures higher than 100 °C (Mauritz and Moore, 2004; Kocha, 2003; Neergat et al., 2003).
6.3
Overview of hybrid inorganic–organic polymer electrolytes
In inorganic–organic polymer electrolytes, inorganic atoms are introduced in the polymer matrix of hybrid PEs. Three classes of inorganic–organic hybrids endowed with promising ionic conductivities will be presented as follows. The first class (Class I) consists of materials prepared by copolymerization of organic macromolecules with metal and non-metal alkoxides and are known as 3D-HION-APE (3-dimensional hybrid inorganic organic networks as polymer electrolytes), discussed by Di Noto et al., 1996a,b, 2002a, 2003a, 2004a; Münchow et al., 2000; Biscazzo et al., 2002; Di Noto, 2002; Vittadello et al., 2002; Di Noto and Zago, 2004. These materials are threedimensional networks where organic macromolecules are bridged by inorganic atoms like Si, Ti, Zr, Al. Their conductivity depends on the doping salts and on the size of the coordination ‘nests’ present in the host material, and is up to 10−5 S cm−1 at room temperature. It is also possible to include in Class I materials nanocomposite systems obtained by dispersing nanopowders of ceramic oxoclusters in perfluorinated polymer electrolytes such as Nafion (Di Noto et al., 2006, 2007, 2008, 2009, 2010a,b; Vittadello et al., 2008; Thayumanasundaram et al., 2010). These hybrid inorganic–organic materials are devised to overcome the drawbacks of conventional perfluorinated proton-conducting PEs. Indeed, with respect to the pristine PE, they are characterized by a higher proton conductivity at higher temperatures (up to ca. 5.9 × 10−2 S cm−1 vs. 3.3 × 10−2 S cm−1 at 115 °C, as reported by Di Noto et al., 2010b). Class I materials are completed by organosiloxanes featuring dangling —SO3H groups. The latter systems are devised as proton-conducting materials and belong to the 3D-HION-APPE family as described by Di Noto and Vittadello (2005) and Di Noto et al. (2005). The second class of inorganic–organic hybrids (Class II) is that consisting of the so-called zeolitic inorganic–organic polymer electrolytes (Z-IOPE), described by Di Noto (1997, 2000), Di Noto et al. (2000, 2001, 2002b, 2003b)
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and Vittadello et al. (2003). This class exhibits the following properties: (a) the organic macromolecules are linked to one another by bridging inorganic clusters; (b) the inorganic clusters are formed by the aggregation of two or more inorganic coordination complexes; and (c) the inorganic clusters can be either positively or negatively charged. These systems are generally prepared by sol → gel → plastic transitions and show a high ionic conductivity (up to 10−5 S cm−1 at room temperature). Hybrid inorganic–organic gels (HGEs) make up the third class of materials (Class III, described by Di Noto et al., 2004b). They are obtained starting from hard and soft metal precursors in a sol → gel process using a lowmolecular weight organic ligand. HGEs can be considered as an upgraded version of Z-IOPEs.
6.3.1 Class I: three-dimensional hybrid inorganic–organic networks as polymer electrolytes (3D-HION-APE) 3D-HION-APEs are reviewed presenting the following examples: • poly[(oligo ethyleneoxide) ethoxysilane] (I) and poly[(oligo ethyleneoxide) ethoxysilane]/(EuCl3)0.67 (II) (Di Noto et al., 1996a,b); • poly[PEG400-alt-DEOS]/(MgCl2)x (6.28 × 10−2 ≤ x ≤ 13.16, PEG = polyethylene glycol, DEOS = diethoxydimethylsilane) (Biscazzo et al., 2002; Di Noto, 2002; Di Noto et al., 2002a; Vittadello et al., 2002); • two poly[(oligo ethyleneglycol) dihydroxytitanate] electrolytic systems (Münchow et al., 2000); • {Zr[(CH2CH2O)8.7]ρ/(LiClO4)z}n (1.80 ≤ ρ ≤ 1.99, 0 ≤ z ≤ 0.90) (Di Noto et al., 2003a) and {Al[(CH2CH2O)8.7]ρ/(LiClO4)z}n (1.85 ≤ ρ ≤ 2.24, 0 ≤ z ≤ 1.06) (Di Noto and Zago, 2004; Di Noto et al., 2004a); • two siloxanic proton conducting membranes (Di Noto and Vittadello, 2005; Di Noto et al., 2005); • proton conducting membranes based on Nafion and [(TiO2)·(WO3)0.148] ‘core–shell’ ceramic oxoclusters (Di Noto et al., 2010b). Poly[(oligo ethyleneoxide) ethoxysilane], i.e. {Si(OEt)1.45 [–O(CH2CH2O)8.7]2.55}n [I] and poly[(oligo ethyleneoxide) ethoxysilane]/ (EuCl3)0.67 i.e. {Si(OEt)1.62 [–O(CH2CH2O)8.7]2.38}n/(EuCl3)0.67 [II] were synthesized by reacting tetraethoxysilane with oligo(ethylene glycol) of molecular weight 400 and oligo(ethylene glycol)400/(EuCl3)0.317, respectively, as shown in Fig. 6.1. The resulting products were very transparent and rubbery; Fourier transform infrared (FT-IR) and Raman studies allowed clarification of the structure of these materials. The latter are cross-linked macromolecular systems where the Si atom is bonded to one –OEt group and to three poly(ethylene oxide) 400 chains. Laser luminescence investigations on [II] showed that Eu3+ ion in the polymer host is located in two different types
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H–O–(CH2–CH2–O)P–H · (EuCl3)w
H–O–(CH2–CH2–O)P–H + w EuCl3 OEt x
Si OEt
+ y H–O–(CH2–CH2–O)P–H · (EuCl3)w
OEt OEt
1. Refluxed at 160–170 °C for 2 h 2. Reduced pressure at 0.1 atm
OEt –(O–CH2–CH2)P–O Si –(O–CH2–CH2)P–O
O–
• (EuCl3)z
(C
H
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CH
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O–(CH2–CH2–O)P–
O)
P
Si O–(CH2–CH2–O)P– OEt n OEt = –O–CH2–CH3 p=8.68 z=0.67
6.1 Preparation procedure of poly[(oligo ethyleneoxide) ethoxysilane]/ (EuCl3)0.67 [II]. Copyright Wiley-VCH Verlag GmbH & Co. Reproduced with permission from Di Noto et al. (1996a).
of sites having a distorted C2v symmetry; the structure of one site is shown in Fig. 6.2. The conductivity of these systems showed a Vogel–Tamman– Fulcher (VTF)-type dependence on temperature. At 70 °C the conductivities of [I] and [II] were 9 × 10−6 S cm−1 and 14.3 × 10−6 S cm−1 respectively. The preparation of poly[PEG400-alt-DEOS]/(MgCl2)x complexes was performed by gradually adding δ-MgCl2 to poly[PEG400-alt-DEOS]. Seven complexes with the formula poly[PEG400-alt-DEOS]/(MgCl2)x with 6.28 × 10−2 ≤ x ≤ 13.16 were prepared by diluting the mother solution 1 : 2n (n = 0, 1, 2, 3, 4, 5, 6) with the pure copolymer poly[PEG400-alt-DEOS]. To obtain a pure polymer host suitable for the preparation of the desired material, the polycondensation reaction between PEG400 and DEOS was performed
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Eu O C Cl
+
Cl(1) O(1) O(2)
Cl(2)
Cl–
O(4)
O(3) (b)
Cl(1)
C2 Cl(2) Eu
O(4) O(3)
O(2)
O(1)
72
6.2 Drawing of the TGT conformational model of a PEO/EuCl3 complex along (a) and parallel (b) to the 72 helix. Copyright Wiley-VCH Verlag GmbH & Co. Reproduced with permission from Di Noto et al. (1996b).
without any catalyst. Analytical data and spectroscopic investigations (1H-, 13C-, and 29Si-NMR (nuclear magnetic resonance), FT-IR) showed that this material is based on host macromolecules whose chains correspond to α-hydro-ω-oligo (oxyethylene) hydroxypoly[oligo (oxyethylene) oxydimethyl sililene], as shown in Fig. 6.3(a). Thus, the copolymer poly[PEG400alt-DEOS] consists of polyethereal moieties linked together by dimethylsililenic units. Solvent-free poly[PEG400-alt-DEOS]/(MgCl2)x electrolytic materials are particularly suitable for the fundamental study of ion–polymer interactions due to the fact that the traces of solvent which can influence the ionic transport properties are completely absent, as described by Ratner and Shriver (1988) and Sun et al. (1996). The polymer poly[PEG400-alt-DEOS] is a highly flexible linear chain based on the PEG400 moieties linked by the dimethylsiloxane bridges. The flexibility is due to the high torsional mobility of the dimethyl sililenic units, which is typical of the polysiloxanic derivatives. Owing to this latter property, it is expected that the typical conformation of PEG400 is maintained in the polyethereal moieties of poly[PEG400-alt-DEOS]. The semiquantitative analysis of the OH and CO vibrational modes provided a deep insight into the structure and the ion–polymer interactions in these systems. Results indicated that: (a) the polyethereal moieties assume
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55
50
45
35
ppm
40
30
25
20 15
10
5
0
–5 –10
ε
HO(CH2CH2O) Si OCH2CH2O(CH2CH2O)CH2CH2O H β β ω 8.671 CH 6.671 ω 3 20.739 ε
(b)
C2
+1
C2
PPD-MgCl2
PPD-[MgCl]+
PPD-Mg2+
6.3 (a) 13C-NMR spectrum of poly[PEG400-alt-DEOS] (PPD) in CDCl3 solution; reproduced from Biscazzo et al. (2002) by permission of Elsevier. (b) Possible coordinations of the Mg2+ cation along the polyethereal chains of PPD; reproduced from Vittadello et al. (2002) by permission of Elsevier.
(a)
85
α
ε
CH3
+2
C2
226
Polymer electrolytes
Fibre period 19.3 Å (a)
Oxygen
Carbon
60° (gauche)
(c) (b)
191.5° (trans)
1Å
6.4 Scheme of the typical 72 helical conformation, factor group D(4π/7), of PEO fragments in polymer electrolytes: (a) lateral view; (b) axial view; and (c) distinction between the trans (T) and the gauche (G) spatial configuration.
a TGT (T = trans, G = gauche) type conformation, as highlighted in Fig. 6.4(c); (b) the doping MgCl2 salt substantially modifies the hydrogen bonding structures of the pure polymer and does not change the chain conformation significantly; (c) the Mg2+ cations are coordinated in a distorted C2v symmetry by the ethereal oxygen atoms of the polyether units. The exact coordination of Mg depends strongly on the ratio between the number of Mg atoms and the number of oxygen atoms, nMg/nO, as shown in Fig. 6.3(b). When nMg/nO ≤ 1.23 × 10−3, the Mg2+ cation is coordinated in a pseudo-tetrahedral coordination geometry (PPD–Mg2+; PPD = poly [PEG400-alt-DEOS]). When nMg/nO = 3.04 × 10−3, the monovalent cationic [MgCl]+ species is present along the polyether chains with a pseudo-trigonal bipyramidal coordination geometry (PPD–[MgCl]+). When 3.04 × 10−3 < nMg/nO ≤ 5.55 × 10−2, the monovalent cationic [MgCl]+ and the neutral MgCl2 species are present with the pseudo-trigonal bipyramidal and pseudooctahedral coordination geometries (PPD–[MgCl]+ and PPD–MgCl2), respectively. When nMg/nO ≤ 1.23 × 10−3, Cl− anions preferentially form hydrogen bonding clusters with the OH groups; as nMg/nO > 1.23 × 10−3, each OH group is hydrogen-bonded to a single Cl− anion. The proposed mecha-
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Hybrid inorganic–organic polymer electrolytes
227
+2 +1 Cl Mg2+ Intra-CH hopping
+2 [Mg-Cl]+ Inter-CH hopping +1
Cl +1
[Mg-Cl]+ Intra-CH hopping
+1 (a)
(b) Oxygen
Carbon
6.5 Models of charge migration in poly[PEG400-alt-DEOS]/(MgCl2)x: (a) inter-chain migration; (b) intra-chain migration.
nisms for the inter- and intra-chain charge migration in these materials are shown in Fig. 6.5. Two 3D-HION-APE composed from poly(oligoethylene glycol) moieties linked together by Ti atoms are prepared by a condensation reaction between anhydrous PEG 400 (or a solution of LiCl in PEG) with Ti(OEt)4 (Münchow et al., 2000). The poly[(oligoethylene glycol) dihydroxytitanate] (I) and poly[(oligoethylene glycol) dihydroxytitanate]/LiCl complex (II) were obtained; the reaction scheme is reported in Fig. 6.6, while the chemical structures of (I) and (II) are reported in Fig. 6.7. In the synthesis reaction two distinct nucleophilic substitutions to Ti(OEt)4 are observed: (a) the terminal hydroxyls of PEG 400 chains coordinate Ti with the subsequent elimination of ethanol and (b) H2O molecules coordinate Ti(IV) atoms with the elimination of ethanol. Vibrational studies show that polyether chains are present in the TGT conformation and that extended hydrogen bond crosslinks occur in these polymers. In polymer (I), these inter-chain interactions occur between hydroxyl groups belonging to titanium atoms and ethereal oxygens, while polymer (II) exhibits both these interactions, together with a second type of hydrogen bonding interaction due to the formation of hydrogen bonding clusters around Cl− ions. The analysis of the
© Woodhead Publishing Limited, 2010
228
Polymer electrolytes
(I) n Ti(OEt)4 + 2n H2O + n HO(CH2CH2O)8.671H
Ti(OH)2O(CH2CH2O)8.671
n
+ 4n EtOH
(II) 1.2 LiCl + HO(CH2CH2O)8.671H
OH(CH2CH2O)8.671H/(LiCl)1.2 Ti(OH)2O(CH2CH2O)8.671
4n Ti(OEt)4 + 8n H2O + 5n HO(CH2CH2O)8.671H/(LiCl)1.2
2 HO(CH2CH2O)8.671H·LiCl
Ti(OH)2O(CH2CH2O)8.671
2
· (LiCl)2 ·(LiCl) · (LiCl)2
+ 16n EtOH n
6.6 Reaction procedures to obtain poly[(oligoethylene glycol) dihydroxytitanate] (I) and poly[(oligoethylene glycol) dihydroxytitanate]/LiCl complex (II); reproduced from Münchow et al. (2000) by permission of Elsevier.
conductivity profiles obtained for these materials evidenced two conductivity regions for polymer (I) and three for polymer (II). All the regions were fitted very well by the empirical VTF equation, suggesting that both polymers conduct ionically by two distinct mechanisms, which are strongly influenced by the segmental motions of the polymer chains. The first conductivity mechanism involves the migration of anionic species from the cathode to the anode due to the presence of reduced titanium species in –[Ti(OH)2–O(CH2–CH2–O)8.671–]n– chains, as outlined in Fig. 6.8. The second conductivity mechanism occurs due to the classical hopping migration of Li+ and Cl− ions. Figure 6.9 reports the VTF graphs for (I) and (II). The thermal stability of these polymers and the possibility of using them to produce thin films are features that make these materials very promising polymer electrolytes. Furthermore, (I) and (II) are rare examples of polymer electrolytes having transition metal atoms in their backbone chains. Eleven network complexes with general formula {Zr[O(CH2CH2O)8.7]ρ / (LiClO4)z}n were prepared by a substitution reaction starting from PEG400, Zr(O(CH2)3CH3)4 and LiClO4 (Di Noto et al., 2003a), as outlined in Fig. 6.10. The resulting materials were transparent and rubbery; 1.80 ≤ ρ ≤ 1.99 and 0 ≤ z ≤ 0.90. FT-Raman analyses showed that: (a) the polyether chains in the bulk materials are present in a TGT conformation; (b) the polymer electrolytes are inorganic–organic network materials with zirconium atoms bonded together by polyether bridges; (c) an unexpected anion-trapping ability toward ClO4− in the hybrid inorganic–organic host matrix is present. The 3D structure of the {Zr[O(CH2CH2O)8.7]ρ/(LiClO4)z}n materials is outlined in Fig. 6.11, while the tetrahedral coordination geometry around the cross-linking Zr atoms is depicted in Fig. 6.12. Impedance spectroscopy studies demonstrated that the proposed systems conduct ionically by a mechanism mediated by the segmental motion and by the concentration of different ionic species distributed in the bulk materials. In particular, two
© Woodhead Publishing Limited, 2010
© Woodhead Publishing Limited, 2010
H
Ti
H
O
O
O
O
O
Ti
H
H
O
= (CH2CH2O)8.671
O
O
H
Ti
H
O
O
O
O
O
O
H
Ti
O
O
H
Ti
H
O
O
O
O H
Ti
H
H
Ti
H
O
O
O
H
O
Ti
O
O
O
O
H
O H
O
O
H
Ti
[Ti(OH)2O(CH2CH2O)8.671]n
H
(I)
H–O
O H
O
O
O
O
O
H
H
Ti
H
O
O
O
Ti
H
Cl–
O
O
Li
H
H
+
O
O
Ti
H
= (CH2CH2O)8.671
O–H = HO(CH2CH2O)8.671H
Ti
H
H
O
Ti
H
H
Ti O
O
O
Li+
H
H
H
Cl–
H
O
O
O
Li+
H O
O
Ti
H
[Ti(OH)2O(CH2CH2O)8.671]
2
HO(CH2CH2O)8.671H·LiCl
2
O
O
Cl– H
O
· (LiCl)2
· (LiCl)2
Ti
O
O
H
H
O
H
Ti
[Ti(OH)2O(CH2CH2O)8·671]
O
H
O
O
O
Li+
Ti
Cl–
H
O
O
O
(II)
O
O
·(LiCl)
H
Ti
H
n
6.7 Structures proposed for poly[(oligoethylene glycol) dihydroxytitanate] (I) and poly[(oligoethylene glycol) dihydroxytitanate]/LiCl complex (II); reproduced from Münchow et al. (2000) by permission of Elsevier.
H
O
Ti
H
H
O
Ti
H
Polymer electrolytes R
R
+ n e–
O (II/III) n– Ti(OH)2
O (IV) 0 Ti(OH)2
O
O – n e–
R
n = 1, 2
R
6.8 Redox mechanism for the conductivity mechanism involving titanium species in –[Ti(OH)2–O(CH2–CH2–O–)8.671–]n– chains of poly[(oligoethylene glycol) dihydroxytitanate] (I) and poly[(oligoethylene glycol) dihydroxytitanate]/LiCl complex (II); reproduced from Münchow et al. (2000) by permission of Elsevier.
(IIC) –8
(IB)
–9
–10 In [σ (T ) / S cm–1]
230
(IIB) –11 (II)' (IIA)
–12 (I) (II) –13
(II)'
–14 (IA)
2.8
2.9
3.0
3.1
3.2
3.3
3.4
T –1 (K–1) × 103
6.9 VTF graphs for poly[(oligoethylene glycol) dihydroxytitanate] (I) and poly[(oligoethylene glycol) dihydroxytitanate]/LiCl complex (II). Conductivity regions of polymers (I) and (II) are indicated; reproduced from Münchow et al. (2000) by permission of Elsevier.
© Woodhead Publishing Limited, 2010
Hybrid inorganic–organic polymer electrolytes (I) PEG400 + y LiClO4
231
PEG400/(LiClO4)y 0≤ y ≤0.49
mα PEG400/(LiClO4)y + m Zr[OCH2CH2CH3]4 4m CH3CH2CH2OH
(II)
{Zr[O(CH2CH2O)8.7]α/(LiClO4)β}m Wash three times with boiling anhydrous toluene {Zr[O(CH2CH2O)8.7]r /(LiClO4)z}n 1.80≤ r ≤1.99
0≤ z ≤0.90
6.10 Reaction steps (I and II) to obtain {Zr[O(CH2CH2O)8.7]ρ/(LiClO4)z}n complexes with 1.80 ≤ ρ ≤ 1.99 and 0 ≤ z ≤ 0.90.
conductivity regions were detected: region I was detected for c1/2 Li ≤ 0.4 −1 1/2 (mol kg−1)1/2, and region II for c1/2 ≥ 0.4 (mol kg ) . Finally, the hybrid Li network with a nLi/nO molar ratio of 0.0223 exhibits a conductivity of ca. 1 × 10−5 S cm−1 at 40 °C, thus suggesting that this solid polymer electrolyte is a good hybrid inorganic–organic ionic conductor. Eleven new hybrid inorganic–organic networks, with the general formula {Al[O(CH2CH2O)8.7]ρ/(LiClO4)z}n, where 1.85 ≤ ρ ≤ 2.24 and 0 ≤ z ≤ 1.06, were prepared by a polycondensation reaction starting from aluminium isopropoxide and polyethylene glycol (PEG400)/(LiClO4)y liquid polymer electrolytes with 0 ≤ y ≤ 0.49 (Di Noto and Zago, 2004; Di Noto et al., 2004a). These materials, which present a glassy rubbery consistency, consist of Al atoms bonded together by PEG400 bridges, as outlined in Fig. 6.13. Thermogravimetric investigations indicated that they are thermally stable up to 260 °C. Medium FT-IR and FT-Raman studies showed that the polyether moieties exhibit a TGT conformation with a helical geometry and detected the presence of Li+. . .ClO4−. . .Al[O(CH2CH2O)8.7]3 neutral species at c1/2 Li ≥ 0.4 (mol kg−1)1/2. Impedance spectroscopy studies, yielding results exemplified in Fig. 6.14, showed that the {Al[O(CH2CH2O)8.7]ρ/(LiClO4)z}n materials conduct ionically by a charge transfer mechanism mainly regulated by the segmental motion and fast ion-hopping processes between equivalent coordination sites distributed along the polyether chains. It was evidenced that in bulk {Al[O(CH2CH2O)8.7]ρ/(LiClO4)z}n complexes the acid aluminium crosslinking sites, which are characterized by two different coordination geometries as reported in Fig. 6.15, exhibit the anion trapping phenomenon
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232
Polymer electrolytes OH Zr
Zr O– +
Zr Zr
OH
Zr
Zr
Zr Zr
O– Zr + Zr
Zr
Zr
Zr OH
O–
Zr
Zr
Zr
Zr
Zr
Zr Zr
O– Zr
Zr
+ Zr
Zr Zr
Zr
Zr
O–
Zr
Zr
Zr
+ Zr
Zr
Zr
Zr
O–
–
O
OH Zr
Zr + Zr
Zr
Zr
Zr
Zr Zr
–
O + Zr
Zr OH
OH
OH
O– Zr
+ Zr
Zr Zr
Zr Zr
Zr
Zr
Zr
Zr Zr
Zr
+ Zr
Zr
Zr Zr
O–
Zr Zr
Zr
+ Zr
Zr
+ Zr
O–
Zr +
+ Zr
OH
Zr+ Zr
Zr
O– Zr + Zr
OH
Zr
Zr Zr
O–
Zr
Zr + Zr
O–
Zr
Zr
OH
Zr Zr
Zr
Zr
Zr
Zr Zr
Zr Zr
Zr
+ Zr
Zr
Zr Zr
Zr
O–
Zr
O–
+ +Zr OH Zr OH Zr+
Zr
OH
Zr
Zr
Zr
Zr Zr
Zr
OH = CIO4–
= Li+
= –O(CH2CH2O)7.671CH2CH2O–
6.11 3D structure of {Zr[(CH2CH2O)8.7]ρ/(LiClO4)z}n materials.
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Hybrid inorganic–organic polymer electrolytes
233
PEG O Zr O
PEG
PEG
O O
PEG
6.12 Tetrahedral coordination geometry around the crosslinking Zr atoms in {Zr[(CH2CH2O)8.7]ρ/(LiClO4)z}n materials. OH Al
Al Al
HO
Al Al
HO Al
O–
HO Al
Al
Al Al HO
H O
H
Al
O–H
Al HO
Al
Al Al
H O Al Al OH
OH
Al
OH O H + Al
H O Al Al Al
Al O H
Al
Al
OH H Al O
HO
Al
Al Al
OH Al
Al HO
OH
Al Al
H O
Al Al
= ClO4–
Al +
HO O
OH Al
H O
HO
Al
Al
Al + Al
HO
Al + O–H
Al
O
O
O H HO
Al
O–H + Al
O–H
H
Al
Al
OH
Al Al
Al
Al
Al+
Al HO
O H
Al Al
Al
HO
Al
OH
+ Al
HO
HO
OH
HO
= Li+
= –O(CH2CH2O)7.671CH2CH2O–
6.13 Structural model proposed for {Al[O(CH2CH2O)8.7]ρ/(LiClO4)z}n materials; reproduced with permission from Di Noto and Zago (2004). Copyright 2004, The Electrochemical Society. © Woodhead Publishing Limited, 2010
234
Polymer electrolytes
18 °C
15 10 5 0 0
50
100 150 200 r´ (Ω cm) × 103
250
1.0 – r˝ (Ω cm) × 106
20
80 °C
0
(a)
– r˝ (Ω cm) × 103
ω
140 120 100 80 60 40 20 0
ω
0.8
18°C 0.6 80°C
0.4
5
300
– r˝ (Ω cm) × 103
– r˝ (Ω cm) × 103
25
0.2
0 0.0
0.5
1.0 1.5 r´ (Ω cm) × 106
30
160 140 120 100 80 60 40 20 0
0.0 (b)
10 15 20 25 r´ (Ω cm) × 103
2.0
50 100 150 r´ (Ω cm) × 103
200
6.14 Nyquist plots for selected {Al[O(CH2CH2O)8.7]ρ/(LiClO4)z}n materials. (a) {Al[O(CH2CH2O)8.7]2.24/(LiClO4)1.06}n and (b) {Al[O(CH2CH2O)8.7]1.89/ (LiClO4)0.01}n. The measurements were carried out from 20 Hz to 1 MHz; the data near the origin of the plots on the left are magnified in the right panels. Reproduced with permission from Di Noto et al. (2004a). Copyright 2004, The Electrochemical Society.
PEG PEG
H
O
O PEG
AI PEG (a)
PEG
O
AI
O
O
O
PEG
O PEG
(b)
6.15 Coordination geometries around the crosslinking Al atoms in {Al[O(CH2CH2O)8.7]ρ/(LiClO4)z}n materials: (a) Four-coordinate Al sites (more abundant); (b) three-coordinate Al sites.
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Hybrid inorganic–organic polymer electrolytes
235
toward perchlorate anions. The best conductivity observed in these materials is 1.66 × 10−5 S cm−1 at 25 °C. 3D-HION-APPE proton conducting materials are presented as an important development to replace the ubiquitous and costly perfluorinated ionomers currently in use. These materials can be synthesized by the copolymerization of metal alkoxides and organic-substituted silicon alkoxides or suitable oligomers. In particular, two new siloxanic proton-conducting membranes are described (Di Noto and Vittadello, 2005; Di Noto et al., 2005): [I] {Si(CH3)3O[Si(CH3)HO]21.26[Si(CH3)((CH2)3SO3H)O]1.8 [Si(CH3)((CH2)3Si(CH3)2O)O]14Si(CH3)3}n; [II] {Si(CH3)3O[Si(CH3)HO]21.26[Si(CH3)((CH2)3SO3H)O]1.8 [Si(CH3)((CH2)3(Si(CH3)2O)w)O]v[Si(CH3)((CH2)3Si(CH3)2O–) O]14−vSi(CH3)3}n with w = 20.31. [I] is synthesized by a two step procedure: at first a precursor is prepared by hydrosilylation of allylsulphonylchloride and allyldimethylchlorosilane using polymethylhydrosiloxane (Fig. 6.16); afterwards, the hydrolysis of this precursor produced a cross-linked silicone polymer where pendant silicone chains are endowed with –SO3H acid groups (Fig. 6.17). [II] is prepared by the polycondensation of the same precursor used in the preparation of [I] with α,ω-dihydroxy(polydimethylsiloxane), obtaining another crosslinked silicone polymer (Fig. 6.18). The two proton-conducting polymer electrolytes [I] and [II] show a remarkable chemical stability and do not decompose until ca. 200 °C. [I] and [II] can be easily transformed into polymer films by sintering processes, and are characterized by Tg values typical of crosslinked silicone materials at −44 and −60 °C, respectively. [I] and [II] have ionic exchange capacities of 0.33 and 0.15 meq g−1, respectively. In [I] and [II] the charge transport takes place owing to an inter-cluster migration of protons through a vehicular mechanism mediated by rotational relaxations of silicone side groups bearing sulphonic acid groups. In this case, the clusters comprise water molecules solvating the sulphonic acid groups. The plot of ln σDC vs. 1/T for [I] and [II] is shown in Fig. 6.19. The σDC at 125 °C is ca. 1.9 × 10−3 S cm−1 and 1.8 × 10−4 S cm−1 for [I] and [II], respectively, at the highest hydration. These values classify these materials as good proton conductors. NMR proton diffusion data on a non-activated [II] membrane showed a two-region behaviour: below 60 °C the sample shows good diffusion properties but at higher temperatures water is quickly released from the sample. A drastic treatment with sulphuric acid reduces the sample into a gummy powder with very high diffusion coefficients all the way up to 110 °C. Taken together, these results are encouraging with respect to the
© Woodhead Publishing Limited, 2010
236
Polymer electrolytes SO3– Na+ POCI3
Na2SO3
Br
SO2CI
(a)
CH3 H3C Si (b)
O
CH3
CH3
Si O
Si
H
CH3
35.42
SO2CI
1/2 CH3
+ Si(CH3)2CI
CH3 1/2
RhCI[(C6H5)3P]3 in 1,2-dimethoxyethane
H3C Si CH3
CH3
CH3
CH3 O
Si H
O
Si x
CH2 H2C CH2 CIO2S
CH3
CH3 O
Si y
CH2 H2C CH2
O
Si z
CH3 (P)
CH3
H3C Si CH3 CI
6.16 Synthesis of the precursor necessary for the preparation of the siloxanic proton-conducting materials [I] and [II]. (a) Preparation of allylsulphonylchloride; (b) hydrosilylation of allylsulphonylchloride and allyldimethylchlorosilane using polymethylhydrosiloxane; reproduced from Di Noto and Vittadello (2005) by permission of Elsevier.
possibility of preparing membranes with performances comparable to those of Nafion. Promising membranes were obtained by doping Nafion with ‘core–shell’ [(TiO2)·(WO3)0.148] ceramic nanoparticles (Di Noto et al., 2010b). This research has been carried out in the framework of a series of efforts aimed at elucidating the effect of ceramic oxoclusters on the structure and the conductivity of hybrid nanocomposite systems (Di Noto et al., 2006, 2007, 2008, 2009, 2010a,b; Vittadello et al., 2008; Thayumanasundaram et al., 2010). The [(TiO2)·(WO3)0.148] nanofiller was obtained by grinding together a dimethylformamide (DMF) suspension of TiO2 and WO3 in the desired ratio (Fig. 6.20). Seven homogeneous membranes with the formula {[Nafion/ [(TiO2)·(WO3)0.148]} where 0 ≤ wt%TiO2 ≤ 15 were prepared by solventcasting. The membranes had a thickness lower than 350 µm, and were stable up to 170 °C. Vibrational spectroscopy investigations allowed us to distin-
© Woodhead Publishing Limited, 2010
Hybrid inorganic–organic polymer electrolytes CH3 n • CH3 Si
CH3
CH3 O
O
Si
O
x
y
CH2
O
Si
CH2
H
CH3
CH3
CH3
Si
CH2 SO2
237
Si
CH2 z CH2 CH2 CH3 Si CH3
CH3
CH3
CI
CI
Hydrolysis in 0.1 M H2SO4 at 60 °C for 12 h
CH3 CH3 Si
CH3
CH3
CH3 O
Si H
O
Si
O
Si
CH2 x
CH2 CH2 SO3H
CH3
CH3
y
O
CH2 CH2
Si z
CH3
CH3 n
CH2 CH3 Si
CH3
O
6.17 Synthesis of the siloxanic proton-conducting polymer electrolyte [I]; reproduced from Di Noto and Vittadello (2005) by permission of Elsevier.
guish hydrophobic and hydrophilic domains at the nano-scale. In particular, it was shown that the fluorocarbon backbone chains are present with two distinct helical conformations, 157 and 103 as outlined in Fig. 6.21; four different species of water domains were also detected. Both groups of vibrational modes are modulated by the amount of [(TiO2)·(WO3)0.148] present in the material. Modulated differential scanning calorimetry (MDSC) analysis revealed that the membranes undergo four thermal transitions at temperatures lower than 300 °C. As the temperature is increased, the transitions are assigned to the melting of small and imperfect fluorocarbon domains in Nafion (100 < T < 150 °C), to the decomposition of the –SO3H groups of Nafion (170 < T < 230 °C) and to the melting of different fluorocarbon domains of Nafion, either not stabilized (230 < T < 270 °C ) or stabilized by the [(TiO2)·(WO3)0.148] nanofiller (T ≈ 300 °C). The temperature of these transitions depends on the membrane composition, as the nanofiller interacts with the sulphonic acid groups of Nafion, improving their thermal stability.
© Woodhead Publishing Limited, 2010
238
Polymer electrolytes CH3 H2O
CI
CI
Si
CH3 CH3 CH3 Si CH3
CH3
CH3 O
Si H
O
O Si CH2 x
CH2 CH2 SO2
Si
Si z
CH3 + Zi• OH
CH3
CH3
Si H
Si
O x
CH2 CH2 CH2 SO3H
H 19.21
Hydrolysis in 0.1 M H2SO4 at 60 °C for 12 h
CH3
CH3
CH3 O
O
Si
CH3
CI
CH3
CH3
O
CH2 y CH2 CH2 CH3 Si
CI
CH3 Si
CH3
CH3
CH3
Si
O y
CH3
CH3 O
CH2 CH2 CH2
z–w
CH3 Si CH3 O
Si O CH2 CH2
Si CH3 CH3 w
CH2 CH3 Si O CH3
n
20.31
6.18 Synthesis of the siloxanic proton-conducting polymer electrolyte [II]; reproduced from Di Noto and Vittadello (2005) by permission of Elsevier.
Dynamic mechanical analysis (DMA) investigations revealed four distinct mechanical relaxation events. α and α′ modes are detected at T > 100 °C and were assigned to the long-range motions of both the backbone and the side chains facilitated by the weakening of electrostatic interactions within the ionic aggregates. β and β′ relaxation events, measured at 35 and 65 °C, were attributed to the 136 → 157 and to the order–disorder conformational transitions occurring in hydrophobic polytetrafluoroethylene (PTFE)-like domains of Nafion, respectively. The electrical response of the membranes was investigated by broadband dielectric spectroscopy (BDS). Figure 6.22 reports the plot of ln σDC vs. 1/T for {Nafion/[(TiO2)·(WO3)0.148]} nanocomposite membranes with 0 ≤ wt%TiO2 ≤ 15. It is highlighted that the nanocomposite materials show a σDC higher by 25–60% with respect to pristine Nafion at 5 < T < 155 °C. σDC reaches a maximum value equal to 5.9 × 10−2 S cm−1 at T = 115 °C for the membrane doped with 5 wt% of TiO2. In the same conditions, the conductivity of pristine Nafion is equal to 3.3 × 10−2 S cm−1. In addition, it is observed that nanocomposite membranes are characterized by a wider Stability Range of Conductivity (SRC) with respect to pristine Nafion. SRC is defined as the temperature range where equation 6.1 holds true: ∂σ DC ≤0 ∂ (1 T )
[6.1]
© Woodhead Publishing Limited, 2010
Hybrid inorganic–organic polymer electrolytes
239
10–2
log (σ / S cm–1)
10–3
10–4
10–5
Nafion 105, 29.78 –wt% of H2O [I] 33.12 –wt% of H2O [II] 27.76 –wt% of H2O fit
10–6 2.6
2.8
3.0
3.2
3.4
3.6
3.8
1/T (K–1) x 10–3
6.19 Temperature dependence of direct current conductivity, σDC, of membranes [I], [II] and Nafion 105 at maximum hydration; reproduced from Di Noto et al. (2005) by permission of Elsevier.
In detail, 0 < SRC < 105 °C for pristine Nafion and 0 < SRC < 145 °C for the {[Nafion/[(TiO2)·(WO3)0.148]} material with 15 wt% of TiO2. A detailed study of the conductivity mechanism has been carried out for other similar materials (Di Noto et al., 2006), and it is suggested that proton migration along interconnecting channels and polar hydrophilic clusters takes place owing to proton exchange processes modulated by the amount and types of interstitial water domains, the density of hybrid nanofiller–(HSO3)–crosslinks, and the segmental motions of the fluorocarbon backbone of Nafion polymer. In turn, this leads to the hypothesis that the exchange of protons between different fluctuating water species domains occurs through hopping processes, and that this mechanism is strongly regulated by the molecular relaxation events in the materials.
6.3.2 Class II: zeolitic inorganic–organic polymer electrolytes (Z-IOPE) Z-IOPEs are hybrid inorganic–organic polymer electrolytes characterized by a structure that can be described as a network of inorganic clusters, © Woodhead Publishing Limited, 2010
240
Polymer electrolytes 1.5 g (≈70%)
0.650 g (≈30%)
DMF
‘Hard’ TiO2
‘Soft’ WO3
Planetary ball mill
5h 500 rpm
with grinding jars of tungsten carbide 0.148 = Dispersion A Mohs Mohs hardness hardness 3 6
molWO3 molTiO2
[(TiO2)(WO3)0.148] nano-clusters Nanometric ‘core–shell’ oxoclusters of a ‘hard’ TiO2 core homogeneously covered by a layer of ‘soft’ WO3.
Ball WO3 milling TiO2 oxocluster oxocluster
The size of ‘core–shell’ oxoclusters ranges from 20 to 60 nm.
6.20 Solid state synthesis of [(TiO2)·(WO3)0.148] ‘core–shell’ nanofiller; reproduced from Di Noto et al. (2010b) by permission of Elsevier.
(a)
(b)
6.21 Axial view of Nafion model compound with different conformations of the backbone: (a) 157 helix; (b) 103 helix.
formed by the aggregation of two or more inorganic coordination complexes either positively or negatively charged, bridged by organic macromolecules. The general structure of Z-IOPEs is shown in Fig. 6.23. Z-IOPEs are prepared starting from two different solutions. The first contains a hard
© Woodhead Publishing Limited, 2010
Hybrid inorganic–organic polymer electrolytes –1.2
241
(II)
–1.3
log (σDC / S cm–1)
–1.4 –1.5 –1.6 –1.7
(I)
wt%TiO2 0 3 5 9 11 13 15
–1.8 –1.9 –2.0 –2.1 –2.2
2.3 2.4 2.5 2.6 2.7 2.8 2.9 3.0 3.1 3.2 3.3 3.4 3.5 3.6 × 10–3 1/T (K–1)
6.22 Dependence of ln σDC on 1/T for {Nafion/[(TiO2)·(WO3)0.148]} nanocomposite membranes with 0 ≤ wt%TiO2 ≤ 15. I and II indicate the conductivity regions. Dotted lines show VTF fitting curves reproduced from Di Noto et al. (2010b) by permission of Elsevier.
n–
n–
n–
n–
= Inorganic cluster n–
n–
n–
= Organic macromolecule
n–
6.23 General structure of Z-IOPEs.
transition metal (Fe, Co) cyanometallate and the organic macromolecule; the second includes a soft metal (Pd, Sn) chloride and the same organic macromolecule. After mixing the two solutions a sol → gel → plastic transition occurs, resulting in the hybrid network. This general synthesis protocol was used to prepare the following Z-IOPEs: [I] [Fe0.922Pd0.663(CN)5.407Cl1.511(C2nH4n+2On+1)K2.826] (Di Noto, 1997); [II] [Fe0.153Pd0.277(CN)0.876Cl0.940(C2nH4n+2On+1)K0.803] (Di Noto, 1997);
© Woodhead Publishing Limited, 2010
242
Polymer electrolytes [III] [Fe0.048Pd0.077(CN)0.103Cl0.44(C2nH4n+2On+1)K0.245] (Di Noto, 1997); [IV] [Co0.110Pd0.273(CN)0.591Cl1.295(C2nH4n+2On+1)K1.01] (Di Noto, 2000); [V] [Fe0.127Sn0.164(CN)0.617Cl0.616(C2nH4n+2On+1)K0.323] (Di Noto et al., 2000); [VI] [Fe0.088Sn0.187(CH3)0.374(CN)0.578Cl0.444(C2nH4n+2On+1)K0.472] (Di Noto et al., 2001, 2002b); [VII] [Fe0.082Pd0.149(CN)0.492Clv(C2nH4n+2On+1)Li0.523] (Di Noto et al., 2003b; Vittadello et al., 2003).
[I], [II], [III] and [IV] were obtained from a first solution composed of K3Fe(CN)6 (or K3Co(CN)6 in the case of [IV]), H2O and PEG600, and a second solution composed of K2PdCl4, H2O and PEG600. By modulating the stoichiometry of the reactants different networks, [I], [II] and [III] were produced. These Z-IOPEs are thermally stable up to ca. 210 °C. The mechanism of the sol → gel → plastic transition was clarified by rheological studies, yielding results exemplified in Fig. 6.24. After mixing the two starting solutions, a direct substitution of the Cl− in PdCl42− by H2O or by the nitrogen atoms of the cyanometallate ligands is likely to predominate, giving rise to a viscoelastic solution according to the reactions shown in Fig. 6.25. Afterwards, H2O is displaced by the hydroxyls groups of PEG600 and the nitrogen atoms of the cyanometallate ligands, giving rise to the transition viscoelastic solution → hard gel, which occurs according to the reactions shown in Fig. 6.26. This reaction mechanism is strongly supported by these observations: (a) if the starting solutions include PEG600 in the absence of water, no gel is formed as the solutions are mixed; (b) the volume of the gel decreases vs. time; at the same time, the gel releases a transparent liquid composed mainly of H2O and traces of PEG600, K, Cl and Pd; (c) the use of water/methanol or water/ethanol solvent mixtures leads to the formation of unstable materials. As a consequence, it is reasonable to assume that the gel-plastic transition occurs mainly owing to a slow, indirect substitution reaction. The conductivity of Z-IOPE [IV] at room temperature is ~3 × 10−5 S cm−1; since the material is very stable and easy to prepare, it appears to be a promising candidate for the development of a new class of electric energy storage systems. The Z-IOPEs [I], [II] and [III], whose structures are reported in Fig. 6.27, show conductivities higher than electrolytic systems based on PEG, poly(ethylene oxide) (PEO) and poly(propylene oxide) (PPO) doped with inorganic salts such as LiClO4, LiCF3SO3, LiSCN, etc. (MacCallum and Vincent, 1989, p285; Scrosati, 1993, p182). In particular, [I] exhibits a conductivity of ≈1 × 10−3 S cm−1 at T = 290 K, which is approximately two orders of magnitude higher than that observed for the electrolyte polymers
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243
(a) 30.3 °C 40.1 °C 49.3 °C 59.7 °C 70.1 °C
G´ (Pa)
1000
100
10
1
0
10
20 Time ×103 (s)
30
(b) 30.3 °C 40.1 °C 49.3 °C 59.7 °C 70.1 °C
G˝ (Pa)
1.5
1.0
0.5
0
10
20 Time ×103 (s)
30
6.24 Typical plots of the elastic G′ (a) and viscous modulus G″ (b) versus time for the sol–gel transition occurring during the preparation of a Z-IOPE at various temperatures; reproduced from Di Noto (2000) by permission of the American Chemical Society.
currently used in electric energy storage devices (Scrosati, 1993, p182; Guyomard and Tarascon, 1994). The temperature dependence of the conductivities was of the VTF type. VTF fitting parameters help us to understand that: (a) [I], [II], and [III] conduct ionically; (b) the increase in the configurational entropy of these inorganic–organic networks is an important factor for the mobility of free ions; and (c) as the concentration of PEG600 in these networks increases, their conductivity decreases. Z-IOPEs [V] and [VI] were obtained applying the general synthesis protocol mentioned above, using SnCl4 and SnCl2(CH3)2 as the soft metal reagent for [V] and [VI], respectively.
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Polymer electrolytes CI
OH–R´–OH
CI
OH–R´–OH
CI
N
M
M
CI
N
–CI– +HO–R´–OH CI
CI
–CI–
(–)
CI
M
M
CI
OH–R´–OH
N
N N CI
CI
–CI– (2–)
M = Pd X = Fe or Co
M
HO–R´–OH = PEG600 CI
CI
N = [X(CN)6]3–
–CI– +H2O CI
(–)
–H2O CI +HO–R´–OH
OH–R´–OH
CI
–CI +HO–R´–OH CI
OH–R´–OH
–H2O +HO–R´–OH
M
N
CI
N –CI–
OH2
CI
CI
OH2
CI
–H2O +HO–R´–OH
N M N –H2O
–H2O
OH2
CI
OH––R´–OH
CI
OH2 M
M CI
CI
M
M
CI
N
CI
–CI– +H2O OH–R´–OH
CI
N
(–)
OH2
–
CI
N (–)
CI M
M CI
N
CI
–CI– +HO–R´–OH
CI
(–)
CI
N N
6.25 Reactions giving rise to a viscoelastic solution in the preparation of a Z-IOPE.
© Woodhead Publishing Limited, 2010
Hybrid inorganic–organic polymer electrolytes CI
N
CI
N +
Pd CI
N
CI
OH–R–OH
N Pd
Pd
N
CI
N
N
N
OH–R–OH Pd
CI
N
HO–––R–HO –2n(CI–)
Pd
CI
N Pd
OH–R–OH
OH–R–OH
N
CI –2CI–
N
N
HO–––R–HO
OH–R–OH
N
Pd
OH–R–OH
245
OH–––R–OH Pd
N
N
OH–––R–OH n
HO–R–OH PEG600 N
N
M = Fe, Co [M(CN)6]3–
N
N
6.26 Typical reactions occurring during the gelification process in the preparation of Z-IOPEs.
[I] NC
[II] 7–
CN CN Fe
NC CN
H
CN
O
Pd CI
NC
CI
R
O H
NC NC NC
CN
CN
•7n K+
CN NC
CN
CN
H
O
O Pd
CI
CN
CI
CI
•2n K+ 4 PEG
Fe NC
CN Fe
R
Pd
Fe
Pd
H
O
O
R
CN
2–
H
H
R
CN CN n
CN
CN
n
[III] –
H
H O
O R
R O
O
O Pd
Pd
H
R
H R
R
CN CN
NC Fe NC
CN
CN
H •n K+ 20 PEG
O H n
6.27 Chemical structure of the Z-IOPEs [I], [II] and [III]. Reproduced from Di Noto (1997) by permission of the Materials Research Society.
© Woodhead Publishing Limited, 2010
246
Polymer electrolytes 5–
N Me OH
C C
Me C
N Sn
HO
Fe
Sn
N C
Me
C
N
OH
N +
* 5K * [PEG600*(KCl)0.254]
C
HO Me
N Me CI
Sn
OH
Me
n
OH
(a)
= (CH2CH2O)12·211 CH2CH2 Temperature (K) 310 320 330 340
300
350 1.5
–13
(I)
In (τ / s)
–9.0
–9.5
(II)
–14
τ1
–15
VTF fit
–16
1.3
p
1.2
–17 In σ (T)
1.1
–18
1.0
VTF fit
–10.0
1.4
d
In (σ (T) / S cm–1)
–8.5
–19 0.9 2.9
(b)
3.0
3.1
3.2
1/T (K–1) ×103
3.3
2.9
3.0
3.1
3.2
3.3
1/T (K–1) ×103
6.28 (a) Structural model proposed for the [VI] Z-IOPE; reproduced from Di Noto et al. (2001) by permission of Elsevier; (b) selected electric properties of [VI]: I) ln[σ(T)] vs. reciprocal absolute temperature; II) ln(τ1) and p as a function of reciprocal absolute temperature. Copyright Wiley-VCH Verlag GmbH & Co. Reproduced with permission from Di Noto et al. (2002b).
The structural hypothesis for [VI] Z-IOPE is shown in Fig. 6.28(a). The TGT conformation of polyether chains is maintained in the final materials. Both [V] and [VI] conduct ionically by a mechanism mainly regulated by the segmental motion of the host material; however, charge migration by ion hopping between equivalent coordination sites in the host network is not to be completely excluded. The conductivities at 25 °C of [V] and [VI] are ca. 3.7 × 10−5 S cm−1 and 4.77 × 10−5 S cm−1, respectively. The electrical properties of [VI] were studied in detail by fitting the data with a universal power law function 6.2; the main results are shown in Fig. 6.28(b):
© Woodhead Publishing Limited, 2010
Hybrid inorganic–organic polymer electrolytes σ ′(ω ) = σ ′( 0 ) ⎡⎣1 + (ωτ1 ) ⎤⎦ p
247 [6.2]
τ1 is a time related to the initial site relaxation rate τ2, while p = τ2/τ* can be expressed in terms of the ratio of the initial back-hopping rate (1/τ*) and the initial site-relaxation rate (1/τ2) in accordance with the jump-relaxation model (see Sections 6.4 and 6.5). [VII] is the first lithium-based Z-IOPE. This material is closely related to its potassium analogues based on K2PdCl4, K3[Fe(CN)6] and PEG600 (Di Noto, 1997). As in the case of the parent potassium Z-IOPEs, the synthesis protocol required the combination of two separate solutions (A and B). Solution A consisted of Li3Fe(CN)6 dissolved in water and PEG600; solution B was obtained by dissolving Li2PdCl4 in water and adding PEG600. The change of cation from potassium to lithium cannot be a simple translation of the chemistry of potassium precursors. Indeed, it is well known that lithium salts are often deliquescent materials and exhibit a high solubility even in organic solvents (Cotton et al., 1999, p92). Therefore, a different behaviour of the reaction mixture used to obtain the lithium Z-IOPE with respect to the potassium Z-IOPE material was expected. First, it was necessary to optimize the preparation of Li2PdCl4 and Li3Fe(CN)6 dry precursors starting from the commercially available compounds. The palladium complex was obtained by reacting LiCl and PdCl2 stoichiometrically in a single step reaction as discussed by Parker and O’Fee (1983): PdCl2 + 2LiCl → Li2PdCl4 The synthesis of the lithium ferricyanide was carried out via a two-step reaction, as described by Chadwick et al. (1985): K3Fe(CN)6 + 3AgNO3 → Ag3Fe(CN)6↓ + 3KNO3 Ag3Fe(CN)6 + 3LiCl → Li3Fe(CN)6 + 3AgCl↓ A possible pathway for the formation of the final product involves the reaction of the two [PdCl4]2− derivatives [PdCl2PEG2] and [PdCl2(H2O)(Fe(CN)6)]3− with each other; their condensation can take place through chloride elimination to produce inorganic–organic clusters that can cross-link, giving rise to a three-dimensional network with the formula [FexPdy(CN)zClv(C2nH4n−2On−1) Lil]. At first, the result of this complexation reaction is the formation of a gel. When the forced release of water under vacuum is completed, a zeolitic polymer electrolyte complex is obtained. The actual final product was a solidplastic material, whose structural model is proposed in Fig. 6.29(a). Elemental analyses and spectroscopic investigations determined that this lithium Z-IOPE is a mixed inorganic–organic network where Fe and Pd are linked by cyanide bridges and Pd atoms are linked by PEG600 bridges. PEG600 assumes the TGT helix conformation. The electrical conduction in this material is mainly due to the displacement of Li+. 1H and 7Li NMR linewidth, spin-lattice
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Polymer electrolytes 2– N R
C
O R
C
N O
Pd
CI
C
H
Fe
N C
O
C
N
2Li+ · [11 PEG600 · 5LiCl]
H
C CI
H
CI
Pd N
R
N
n
(a) –2 I
8
–3 II
6
–4
–5
4 fion s DC
–6
log (sDC / S m–1)
log (fR/Hz)
fseg
2 –7 0 –8
3.0
(b)
3.2
3.4
3.6
3.8
4.0
1/T (K–1) × 10–3
6.29 (a) Structural model proposed for the Z-IOPE [VII]; (b) dependence on temperature of fion, fseg and σDC for Z-IOPE [VII]; reproduced from Vittadello et al. (2003) by permission of Elsevier. fion, fseg and σDC are the frequency of the ion-mode relaxation, the frequency of the relaxation event attributed to the segmental motion and the conductivity of the material, respectively.
relaxation and pulsed field gradient diffusion measurements suggest that the lithium ion transport is correlated with the mobility of the polymer, as in the case of ‘conventional’ polymer electrolytes. The electrical spectra for frequencies higher than 15 kHz evidenced the presence of relaxation events associated to local ion motion dynamics and long-range diffusion. Results are shown
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Hybrid inorganic–organic polymer electrolytes
249
in Fig. 6.29(b). These two phenomena were interpreted in terms of: (a) ion hopping processes; (b) site relaxations; and (c) host medium reorganization processes. Taken together, these studies demonstrated that at temperatures higher than 35 °C, this Z-IOPE conducts ionically by charge transfer mechanisms mainly regulated by ion hopping between equivalent polyether coordination sites followed by correlated host medium reorganizations. Finally, a conductivity of 5.3 × 10−5 S cm−1 at 35 °C classifies this hybrid inorganic–organic network as a good lithium ion conductor.
6.3.3 Class III: hybrid gel electrolytes (HGEs) The HGEs are prepared through a sol → gel process, starting from a combination of hard and soft metal precursors; a low molecular weight organic ligand is used instead of a macromolecule such as PEG600. Two products obtained with glycerol as the molecular ligand are described as examples of this class of materials. Glycerol is a very well-known glass former (Angell, 1997) that has been thoroughly investigated in BDS (Lunkenheimer et al., 2000) and in thermodynamic studies (Ito et al., 1999; Martinez and Angell, 2001). On the basis of the ‘strong–fragile’ classification scheme (Angell, 1997; Ito et al., 1999; Martinez and Angell, 2001), glycerol is considered to be a strong glass former and therefore is characterized by a molecular relaxation that increases with an almost VTF-like (simple activated) profile (Lunkenheimer et al., 2000). The analysis of the frequency-dependent dielectric loss of glycerol reveals that it is a ‘Type A’ glass former (Angell, 1997; Kudlik et al., 1999), since it exhibits a well-defined α-relaxation followed by an excess wing, without any slow β-relaxation as in the case of ‘type B’ glass formers (Kudlik et al., 1999). Therefore, glycerol is a suitable organic molecule to investigate the effect of inorganic clusters in an electrolytic gel structure. Glycerol also represents a promising precursor for the preparation of gel lithium ion conductors able to operate at low temperatures owing to the absence of thermal transitions over a wide temperature range (Angell, 1997; Kudlik et al., 1999). Two systems are presented: [I] [FexPdy(CN)zClv(C3H8O3)Lil] (Di Noto et al., 2004b); [II] [FexSny(CH3)2y(CN)zClv(C3H8O3)Lil] (Di Noto et al., 2004b). The synthesis of these new HGE electrolytic complexes required the combination of two solutions (A and B). [I] was prepared by mixing solutions A and B under an argon atmosphere. Solutions A and B consisted of glycerol dissolving Li2PdCl4 and Li3Fe(CN)6, respectively. Solutions A and B were mixed, and a gel was obtained quickly. The latter expelled a small amount of glycerol over 1 week. The resulting material was filtered under an argon atmosphere and left under vacuum for 10 days. The final product
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Polymer electrolytes a Li3Fe(CN)6 + b glycerol
a K + b glycerol with K = Li2PdCI4; Sn(CH3)2CI2 A
B
+
sol–gel transition
[FexMy(CN)zCIv(CH2OHCHOHCH2OH)Li5]gel M = Pd; Sn(CH3)2
6.30 Synthesis of hybrid inorganic–organic gels based on Li2PdCl4 (I) or Sn(CH3)2Cl2 (II), glycerol and Li3Fe(CN)6; reproduced from Di Noto et al. (2004b) by permission of the American Chemical Society.
(a)
(b)
6.31 Morphology of HGEs determined by scanning electron microscopy (SEM): (a) HGE [I]; (b) HGE [II]; reproduced from Di Noto et al. (2004b) by permission of the American Chemical Society.
was paste-like and dark-brown with a greenish cast. [II] was synthesized by mixing A and B solutions in open air. Solutions A and B contain glycerol dissolving (CH3)2SnCl2 and Li3Fe(CN)6, respectively. As A and B were stirred after the mixing, a yellow-orange gel was immediately obtained. The material was filtered and left under vacuum for 10 days, yielding a yelloworange gel with a slightly green cast. A summary of the synthesis procedure is shown in Fig. 6.30. It is expected that the sol → gel transition occurring in both HGE systems takes place with the same mechanism outlined for conventional Z-IOPEs (Di Noto et al., 2001, 2002b, 2003b). The morphology of [I] and [II] is reported in Fig. 6.31. © Woodhead Publishing Limited, 2010
Hybrid inorganic–organic polymer electrolytes
251
Spectroscopic investigations concluded that these materials consist of mixed inorganic–organic networks containing cavities flooded with glycerol glass-former lithium electrolyte materials. The hybrid inorganic–organic networks in [I] are constituted of Pd atoms bound to Fe by CN bridges, and Pd atoms linked together by glycerol bridges. Likewise, the networks in [II] are composed of Sn atoms bound to Fe through CN bridges and Sn atoms linked together by glycerol. The structural models for [I] and [II] HGEs are shown in Fig. 6.32. The glycerol molecules in lithium electrolyte glass formers
(–) N C
CN
CN Fe
O
H
HO
HO
H
OH
CN
CN
O
Pd
∙ Li+ OH
HO OH
N C
OH
C N
HO OH
OH OH
H
NC
OH
O H
HO
Fe
O
O
H
CN
O
CN
CI
Pd
Pd
OH
OH
CN 244 ∙ HOCH2CHOHCH2OH 12∙LiCI
C N
(a) HO
HO OH OH
OH
n
HOCH2CHOHCH2OH
OH =
CH3
HO Sn
N C
CH3 HO HO
Sn
OH
OH
CN
HO
CN
CN
H3C
CN CH3
Sn
CH3 CH3
CN Fe
H3C
CN
CN
OH
CN H O
O H
OH CN
OH
Sn C N
(b)
CI
CI
O
5 LiCI 272 ∙ HOCH2CHOHCH2OH
–CH3
OH OH
n
OH
H O HO
O Sn
H O
HO
CH3
OH
CI
Fe
N C
HO
HO
OH
OH
O H
OH
HO
OH
HO HO
OH HO
CI
O H
H3C
HO
OH =
HOCH2CHOHCH2OH
6.32 (a) Structural model proposed for HGE [I]; (b) structural model proposed for HGE [II] reproduced from Di Noto et al. (2004b) by permission of the American Chemical Society.
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Polymer electrolytes
αα conformer
αγ conformer
γγ conformer
6.33 Typical conformers of glycerol molecules.
(a)
HGE[I]
HGE[II]
100
50 80 40
Permittivity
20
Permittivity
80 °C 80 °C 30
60 40 20
10 –60 °C
0
0 100
102
104
106
108
100
1010
–60 °C 102
Frequency (Hz)
104
106
108
1010
Frequency (Hz)
(b) 20
40 –60 °C
–60 °C 30
80 °C
Permittivity
Permittivity
15
10
5
0 100
80 °C
20
10
0 102
104
106
Frequency (Hz)
108
1010
100
102
104
106
108
Frequency (Hz)
6.34 Imaginary component of the dielectric spectra of HGEs: (a) dielectric loss spectra and (b) difference dielectric loss spectra obtained by subtracting the spectral contribution of electrode polarization and σDC from dielectric spectra shown in (a). Measurements were carried out from −60 to +80 °C in 10 °C increments; reproduced from Di Noto et al. (2004b) by permission of the American Chemical Society.
© Woodhead Publishing Limited, 2010
1010
Hybrid inorganic–organic polymer electrolytes
253 10–5
0.16
10–6
0.14
10–7 I II
10–8 10–9
0.10
10–10
0.08
D (cm2 s–1)
λ (nm)
0.12
10–11 0.06
I II
10–12
0.04
10–13
0.02
10–14 220
240
260
280
300
320
340
360
Temperature (K)
6.35 Temperature dependence of migration step length λ and diffusion coefficient D in HGE[I] and HGE[II] reproduced from Di Noto et al. (2004b) by permission of the American Chemical Society.
flooding the HGEs’ bulk cavities are predominantly a mixture of αα, αγ and γγ conformers, depicted in Fig. 6.33. Variable-temperature 1H and 7Li NMR linewidth spin-lattice relaxation and pulsed field gradient diffusion measurements suggested that lithium ion transport in the HGEs differs from that in common polyether-based PEs. Indeed, charge transport in HGEs occurs via short-range motions within glycerol solvation clusters. BDS investigations (Fig. 6.34) were in accordance with NMR results and revealed that the conductivity of the two HGEs is modulated by the α and the slow β mode relaxations of glycerol glass-forming molecules. Furthermore, accurate analyses of the temperature dependence of the mode relaxation parameters and of σDC values indicated that ion diffusion in electrolytes flooding the cavities of the HGEs takes place over distances in the order of magnitude of glycerol intermolecular hydrogen bond lengths, as shown in Fig. 6.35. These results permit the conclusion that the two HGEs probably conduct ionically through a charge migration mechanism, which involves short-range exchange of lithium ions between glycerol coordination cages, as shown in Fig. 6.36. The lithium coordination cages are generated owing to intermolecular hydrogen bonds between glycerol molecules forming the solvation layer of the hybrid inorganic–organic network backbone. Finally, [I] and [II] HGEs exhibit room temperature conductivities equal to 5.0 ×
© Woodhead Publishing Limited, 2010
254
Polymer electrolytes GC1
GC2
GC1Li+ + GC2 ↔ GC1 + GC2Li+
6.36 Model of the exchange of lithium ions between glycerol coordination cages.
10−5 S cm−1 and 8.5 × 10−5 S cm−1, respectively. Thus, HGEs can be classified as good lithium electrolyte materials.
6.4
Methods
6.4.1 Structural characterization of hybrid inorganic–organic polymer electrolytes The structural characterization of hybrid inorganic–organic polymer electrolytes and the study of their conductivity requires a holistic approach, integrating data obtained by several techniques. A useful guiding principle for a fruitful investigation of these materials is to build up knowledge of corresponding conventional polymer electrolytes seen as the components of a modular structure. Typically, the structural study of the materials starts from the chemical analysis, carried out by elemental microanalysis and inductively coupled plasma atomic emission spectroscopy (ICP-AES). On the basis of the compositional information it is possible to calculate the ratio between the organic molecules (e.g. PEG and glycerol) and the inorganic component. The comparison between the expected and the experimental composition of the investigated material is helpful to understand the mechanism of the chemical reactions involved in the synthesis. A viscosimetric study in terms of G′(ω) and G″(ω) is useful to monitor the gel formation in Z-IOPEs and HGEs, as shown in Fig. 6.37. Scanning electron microscopy (SEM) and energy-dispersive X-ray spectroscopy (EDX) are used to confirm the morphological and compositional homogeneity of the synthesized materials. Far-infrared (FIR) spectroscopy provides
© Woodhead Publishing Limited, 2010
1
100
0
80
–1
60
–2
40
δ (°)
log (G′/Pa); log (G″/Pa)
(a)
log (G′/Pa) log (G″/Pa) δ (°)
–3
20 0
–4 –1.0
–0.5
0.0
0.5
1.0
1.5
log (ω / rad s–1) 30
3
25
2
20 1 15 0
log (G′/Pa) log (G″/Pa) δ (°)
–1
δ (°)
log (G′/Pa); log (G″/Pa)
(b)
10 5 0
–2 –1.0
–0.5
0.0 0.5 log (ω / rad s–1)
1.0
1.5
log (G′/Pa); log (G″/Pa)
(c) 3.0 2.5 2.0 log (G′/Pa) log (G″/Pa)
1.5 1.0 0.5 0.0 –1
0
1 log (τ / Pa)
2
3
6.37 G′, G″ and δ (deg) vs. frequency plots measured after a reaction time of 35 min (a) or 6 h (b). (c) G′, G″ as functions of the logarithm of the stress (τ/Pa) measured after a reaction time of 6 h. The time when the solutions A and B leading to the Z-IOPE were mixed was taken as t = 0; reproduced from Di Noto (2000) by permission of the American Chemical Society.
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Polymer electrolytes
a qualitative indication about the metal–ligand interactions in the system. The compositional data and the metal–ligand vibrations allow a preliminary structural hypothesis to be made on the inorganic clusters or on the metal/ semi-metal atoms distributed in the material and the organic molecules bridging the inorganic components. Detailed structural information is obtained from medium infrared (MIR) and Raman spectra. Vibrational modes are identified by a correlative analysis based on the literature. The full assignment of the spectra is carried out by identifying the spectral features of the organic component and those of the inorganic component. It is particularly useful to determine the conformation of the macromolecules (e.g. PEG) which are included in the structure of the hybrid inorganic– organic polymer electroytes. An example of FT-IR absorption spectra of a Z-IOPE is shown in Fig. 6.38. In general, it is observed that the macromo(a)
1.2
Absorbance
1.0 0.8 0.6 0.4 0.2 0.0 500
1000
1500
2000
2500
3000
3500
4000
Wavenumber (cm–1)
(b)
Absorbance
1.2 1.1 1.0 0.9 0.8 0.7
100
200
300
400
500
600
Wavenumber (cm–1)
6.38 FT-IR absorption spectra of the Z-IOPE [VII]: (a) MIR; (b) FIR. Reproduced from Di Noto et al. (2003b) by permission of Elsevier.
© Woodhead Publishing Limited, 2010
1000
Residuals
Residuals
Hybrid inorganic–organic polymer electrolytes
0
–1000 × 10–6
0.004
0.000
–0.004
60 × 10–3
HGE [I]
Absorbance (a.u.)
Absorbance (a.u.)
30 ×10 –3
20
10 Current
257
HGE [II]
40
20
pesk
0
0 1900
2000
2100
2200
Wavenumber (cm–1)
2300
2000
2100
2200
Wavenumber (cm–1)
6.39 Decomposition by Gaussian functions of the MIR FT-IR spectral range from 1900 to 2500 cm−1 of HGE [I] and HGE [II]; reproduced from Di Noto et al. (2004b) by permission of the American Chemical Society.
lecular conformation of PEG in conventional PEs such as PEG/(MX)n and in the hybrids overviewed in this chapter is the same. PEG is characterized by a TGT conformation that can be interpreted using the symmetry classes of the D(4π/7) group (Di Noto et al., 1996b). PTFE fluorocarbon chains of Nafion show either a 157 or a 103 helical conformation (Fig. 6.21). The relative contribution of the relevant bands in the MIR spectra is determined by Gaussian decomposition of the appropriate spectral regions (Di Noto et al., 2006, 2007, 2008, 2010b; Vittadello et al., 2008). In the case of the Z-IOPE and HGE materials the structural hypothesis is refined using a semi-quantitative vibrational analysis carried out on the stretching CN and the stretching OH vibrations. The distinctions between bridging and terminal cyanide groups (CNb and CNt, respectively) allows the determination of the structural connectivity of the inorganic clusters, as shown in Fig. 6.39 and Table 6.1. The distinction between bridging hydroxyl groups (OHb) and hydroxyl groups involved in hydrogen bonding cages with other hydroxyl groups (OHHy) allows us to determine the proportion of organic moieties directly linked to the inorganic clusters. In the case of 3D-HION-APE materials and Nafion-based composites, the MIR and Raman spectra are used not only to confirm the structural hypotheses but also to determine ion–ion and ion–polymer interactions by using difference spectroscopy techniques; an © Woodhead Publishing Limited, 2010
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Polymer electrolytes
Table 6.1 Band parameters of the ν(CN) FT-IR vibrational modes for HGE [I] and [II]. Reproduced from Di Noto et al. (2004b) by permission of the American Chemical Society HGE
Band
[I]
ν(–CN)b ν(–CN)t
[II]
ν(–CN)b ν(–CN)t
νi (cm−1)
{ { { {
2218 2185 2158 2126 2063 2025 2144 2103 2073 2038
AiCN a
fwhmia
0.45 ± 0.01 0.21 ± 0.01 0.05 ± 0.01 0.92 ± 0.02 1.97 ± 0.03 0.09 ± 0.01 0.50 ± 0.02 1.02 ± 0.70 2.3 ± 0.3 0.32 ± 0.07
30.4 28.1 17.2 46.4 63.6 24.3 23.3 33.4 34.7 29.6
± ± ± ± ± ± ± ± ± ±
0.5 1.7 2.0 0.5 1.7 2.0 0.7 4.0 3 3
RiCN (%)b
RiCN c,d
RteoCN
12.19 5.69 1.35 24.93 53.39 2.44 12.08 24.63 55.56 7.73
RbCN = 0.19
≈1/6
RtCN = 0.81
≈5/6
RbCN = 0.37
≈2/6
RtCN = 0.63
≈4/6
a
Ai and fwhmi are the band area and the full-width at half-maximum of the peak centred at νi, respectively. b RiCN (%) = (AiCN × 100/ΣAiCN). c RbCN = (ΣAib/ΣAiCN); Aib is the band area of bridging CN groups. d RtCN = (ΣAit/ΣAiCN); Ait is the band area of the terminal CN groups.
example is shown in Fig. 6.40 (Di Noto et al., 2003a). The application of these methods allows the salt–polymer interactions to be determined as the concentration of salt is varied. Further structural indications are obtained by thermal analysis (thermogravimetric (TG) and MDSC). The thermal stability is estimated, the Tg and the phase transitions are measured. The mechanical relaxations of the structure can be measured by dynamic mechanical analysis (DMA). Solid state NMR is particularly useful to determine the structural characteristics of siloxanic systems using 29Si and to probe the environment of Li ions using 7Li. The 1H and 13C spectra are useful to determine the correlation between the Li ions and the organic host on the basis of the spin-lattice relaxation time T1 and the lineshape.
6.40 (a) Difference FT-Raman spectra of {Zr[(CH2CH2O)8.7]ρ/(LiClO4)z}n complexes with 1.8 ≤ ρ ≤ 1.99 and 0 ≤ z ≤ 0.90. The difference spectra were determined by subtracting the spectrum of the pristine hybrid inorganic organic network (cLi = 0) from the spectra of the {Zr[(CH2CH2O)8.7]ρ/(LiClO4)z}n complexes. Dotted line shows the spectrum of pure LiClO4 salt; (b) decomposition by Gaussian functions of difference FT-Raman spectrum of the {Zr[(CH2CH2O)8.7]1.99/(LiClO4)0.90} n complex. The inset shows the dependence of ip% on the lithium concentration, cLi1/2 (mol kg−1)1/2. ip% is the percentage of the ClO4− anions involved in the formation of ion pairs. ip% = A940 × 100/(A940 / A930), where A940 and A930 are the band areas of the peaks at 940 and 930 cm−1, respectively. Reproduced from Di Noto et al. (2003a) by permission from Elsevier.
© Woodhead Publishing Limited, 2010
Intensity (a.u.)
Hybrid inorganic–organic polymer electrolytes
cLi mol kg–1 0.92104 0.41700 0.21758 0.10706 0.05807 0.03055 0.01484
LiCIO4
200
259
400
600
(a)
1200 800 1000 Wavenumber (cm–1)
1.0
1400
1600
25
ip (%)
20
Intensity (a.u.)
0.8
15 10 5
0.6
0 0.2
0.4
0.4 0.6 0.8 cLi½ (mol½ kg–½)
1.0
0.2
0.0 200 (b)
400
600
800 1000 1200 Wavenumber (cm–1)
© Woodhead Publishing Limited, 2010
1400
1600
260
Polymer electrolytes
6.4.2 Conductivity studies on hybrid inorganic–organic polymer electrolytes Conductivity is studied mainly by BDS. This technique is better known as impedance spectroscopy at frequencies lower than 1 MHz. Impedance spectroscopy yields Nyquist plots (−Z″(ω) vs. Z′(ω)), which can be analysed in terms of an equivalent circuit (EC). Figure 6.41 shows a typical example of these plots. Typically, the Nyquist plot of a polymer electrolyte at a given temperature and salt concentration can be simulated by a parallel resistorcapacitor (RC) in series with a constant phase element, as shown in Fig. 6.42.
ω
4000
–Z˝ (ohm)
3000
21.3 °C 31.6 °C 41.0 °C 51.8 °C 61.3 °C 70.7 °C 79.9 °C
2000
1000
0 2000 3000 Z´ (ohm)
1000
4000
6.41 Nyquist plots obtained for the Z-IOPE [VI] at various temperatures and in the frequency ranges from 20 Hz to 1 MHz. Copyright Wiley-VCH Verlag GmbH & Co. Reproduced with permission from Di Noto et al. (2002b).
Zb CPE Cb
6.42 Equivalent circuit used to model Nyquist plots shown in Fig. 6.41.
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Hybrid inorganic–organic polymer electrolytes
261
The capacitance and the constant phase element can be distinguished by using two distinct constant-phase elements (CPE) in the simulation and by determining the corresponding value of nCPE. The parallel RC is represented by a semicircle at higher frequencies, while the Warburg element appears as a linear branch at lower frequencies. The value of R determined by EC analysis is used to calculate the direct current conductivity (σDC = 1/ρ, where ρ is the resistivity of the material). A VTF equation (equation 6.3) can be used to fit the plot of σDC vs. 1/T: σDC (T ) = Aσ ⋅T
−
1 2
⋅ exp [ − EVTF R (T − T0 )]
[6.3]
The VTF equation models a viscous material whose conductivity is regulated by segmental motion. More than one VTF branch is observed in some samples in different temperature regions (Münchov et al., 2000; Di Noto et al., 2010b). The VTF equation has three parameters: Aσ is a pre-exponential factor proportional to the concentration of carriers, EVTF is a pseudoactivation energy for the conduction and T0 is the ideal thermodynamic glass transition temperature [(Tg − 55) ≤ T0 ≤ (Tg − 40)] (Di Noto, 2002). These parameters can be analysed vs. the molar ratio between the concentration of cations and the concentration of oxygen atoms in the backbone or vs. the square root of the concentration of cations to identify regions with different conductivity mechanisms. It was found that the σDC value determined at low frequencies (below 10 kHz) from EC analysis matches the real conductivity at zero frequency σ′(0) determined above 10 kHz with the universal power law (UPL, Equation 6.2). Figure 6.43 shows a typical analysis carried out on the spectra of σ′(ω) in the framework of the UPL. The result σDC = σ′(0) is important since the EC analysis fails to predict σ′(ω) at high frequencies (above 10 kHz). The other two parameters of the UPL are also interesting. τ1 is a time related to the initial site relaxation rate τ2, while p = τ2/τ* can be expressed in terms of the ratio of the initial backhopping rate (1/τ*) and the initial site-relaxation rate (1/τ2) following the jump relaxation model (see Section 6.5). Depending on the value of p it is possible to estimate the effectiveness of forward hopping. On the basis of the value shown by p (either larger or smaller than 1) and knowing what ionic species are present in the system it is possible to distinguish between mostly cationic or mostly anionic transport mechanisms, within different materials in different temperature regions. In the case of Z-IOPEs it was shown that p is decreased as T is raised, as shown in the right panel of Fig. 6.28(b) (Di Noto et al., 2002b). This evidence implies that as T is raised the initial site relaxation time τ2 decreases more quickly than the initial backhopping time τ*. The ionic species are determined by plotting the equivalent conductivity ⌳ vs. the square root of the cationic concentration, as shown in Fig. 6.44. By extrapolating ⌳ to zero concentration it is also
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262
Polymer electrolytes [II]
400
80° C
300 200 100 18° C
0 2
(a)
3
5
4
σ´ (ω) (S cm–1) × 10–6
σ´ (ω) (S cm–1) × 10–6
[I] 16
80° C
14 12 10 8 6 4 2
18° C
0
6
3
2
4
5
6
log (f/Hz)
log (f/Hz) 2.4
1.4
2.2
σ´ (ω) / σ´(0)
σ´ (ω) / σ´(0)
2.6 1.6 1.2 1.0 0.8 0.6 0.4
1.6 1.4 1.0
2
3
4
5
6
2
3
In[σ´ (ω) / σ´(0)–1]
0
18 °C
–1 –2 –3 –4 –5 –6
80 °C 5.4
5.5
5.6
5.7
4
5
6
log (f/Hz)
log (f/Hz)
(b) In[σ´ (ω) / σ´(0)–1]
1.8
1.2
0.2
(c)
2.0
5.8
5.9
1 18 °C
0 –1 –2 –3 –4
80 °C 5.2
6.0
5.4
log (f/Hz)
5.6
5.8
6.0
log (f/Hz)
6.43 Typical analysis of the real component of the complex conductivity σ′(ω) in the framework of the UPL. (a) Plots of the raw data fitted with the equivalent circuit shown in Fig. 6.42; (b) plots of the normalized figure σ′(ω)/σ′(0); and (c) plots of ln[σ′(ω)/σ′(0)-1]; the angular coefficient is the exponent p of the UPL. [I] and [II] are the materials {Al[O(CH2CH2O)8.7]2.24/(LiClO4)1.06} and {Al[O(CH2CH2O)8.7]1.89/ (LiClO4)0.01}n, respectively. Reproduced with permission from Di Noto et al. (2004a). Copyright 2004, The Electrochemical Society.
possible to determine ⌳0, which is correlated to the diffusion coefficient. A typical ⌳0 vs. T example is detailed on Fig. 6.45. The UPL equation can be generalized in the GUPL equation 6.4 (Di Noto, 2002): ⎡ p ⎤ σ ′(ω ) = σ ′( 0 ) ⎢1 + ∑ fi ( τ1,i ω ) i ⎥ ; ∑ fi = 1 ⎣ i =1 ⎦ i =1 N
N
© Woodhead Publishing Limited, 2010
[6.4]
Hybrid inorganic–organic polymer electrolytes 293.15 K 303.15 313.15 323.15 333.15 343.15 353.15 fit
1.0 Λ [S kg (cm mol)–1] × 10–3
263
0.8 0.6 0.4 0.2 0.0 0.2
0.4
0.6
1.0
0.8
cLi½ (mol/kg)½
6.44 Equivalent conductivity (⌳) as a function of salt concentration in {Al[O(CH2CH2O)8.7]r/(LiClO4)z}n complexes. Reproduced with permission from Di Noto et al. (2004a). Copyright 2004, The Electrochemical Society.
–3.2 log (Λo) fit
log [Λo / S kg (cm mol)–1]
–3.4 –3.6 –3.8 –4.0 –4.2 –4.4
2.9
3.0
3.1 3.2 1/T (K–1) × –3
3.3
3.4
6.45 Equivalent conductivity at infinite dilution (⌳0) vs. the reciprocal of temperature in {Al[O(CH2CH2O)8.7]r/(LiClO4)z}n complexes. Reproduced with permission from Di Noto et al. (2004a). Copyright 2004, The Electrochemical Society.
© Woodhead Publishing Limited, 2010
264
Polymer electrolytes 0
6 –2
log e /e 0
eel e´
2 0 –2
e˝
Dielectric relaxation
Electrode polarization –3
0
tion 3 log f [Hz]
6
tseg 9
log σ [S/m]
DC conduction 4
Electrode polarization
–4
tel DC conduction
–6
s´
Dielectric relaxation
sDC s˝
–8 –3
0
3 log f [Hz]
6
9
6.46 Electrical response of materials in terms of dielectric (left) and conductive spectra (right).
The GUPL equation includes the possibility of multiple relaxations characterized by a different fractional contribution fi. The application of BDS methods up to 1 GHz allows us to measure directly the ionic and segmental relaxation phenomena. These measurements yield the complex permittivity ε*(ω) and the complex conductivity σ*(ω). The latter two functions are complementary, as shown in equations 6.5 (Furukawa et al., 1997; Di Noto, 2002): σ′(ω) = ω · ε″(ω) σ″(ω) = ω · ε′(ω)
[6.5]
ε*(ω) and σ*(ω) can be used to emphasize different information, as highlighted in Fig. 6.46. ε*(ω) is particularly suitable to visualize the ionic and segmental relaxation phenomena (τion and τseg) after subtracting the contribution arising from σDC. By comparing σDC, τion and τseg as a function of the reciprocal temperature 1/T it is possible to determine whether the mechanism dominating the conductivity mechanism is the ionic phenomenon or the segmental motion (Vittadello et al., 2003). The measurement of ε″(ω) by BDS also provides a method to distinguish between different types of molecular motions in the polymer host, i.e. the segmental mode, the local mode β and the normal mode n (taking place between end groups), as shown in Fig. 6.47. The distinction is possible on the basis of the different dependence of the logarithm of the relaxation time (τα, τβ and τn) as a function of the reciprocal temperature: ln(τβ) is linear, while ln(τα) and ln(τn) show a Vogel–Tamman–Fulcher–Hesse (VTFH) behaviour, shown in equation 6.6 (Furukawa et al., 1997; Di Noto, 2002): ln ( τ ) = ln ( τ 0 ) −
A (T − T0 )
© Woodhead Publishing Limited, 2010
[6.6]
Hybrid inorganic–organic polymer electrolytes Segmental and ionic motions
10
265
0
8
–2
fseg
O 6
X–
log f (Hz)
O
O
X
O
O
–
M+ O
–4 sDC
4
–6
2
–8
0
–10
–2
–12
Tg
O
–100
log s (S/m)
O M+
–50
0
50
100
Temperature (°C)
1010
10–4
108
10–6 fseg sDC
106
10–8
104
10–10
102
10–12
100
10–14 Tg
10–16
10–2 0
1
2
3
4
sdc (S m–1)
fseg (Hz)
6.47 Example of the dependence vs. temperature of σDC and segmental mode.
5
–3 –1
1/T (10 K )
6.48 VTFH simulations of conductivity vs. T profiles. The conduction mechanism is regulated by the segmental motion.
A good fit of data with a VTFH equation indicates that the conductivity mechanism is regulated by segmental motion, as exemplified in Fig. 6.48. The segmental mode and the n normal mode can be distinguished on the basis of their position on the frequency axis; indeed, mode n is characterized
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Polymer electrolytes
by lower frequencies. The measurements of the molecular motions in siloxanic proton conductors was particularly insightful in the identification of the role of the sulphonic groups. The latter give rise to the β mode and promote proton mobility among water domains. Solid state NMR can be used to determine: (a) the diffusion coefficients of protons and metal cations and; (b) the activation energies for conduction. As a consequence, this technique provides an independent method to validate the results of the electrical analysis (Di Noto et al., 2005).
6.5
The real component of the conductivity spectra in the framework of the jump relaxation model and polymer segmental motion
The discussion of the conductivity mechanism of hybrid inorganic–organic materials can benefit by taking into account the real component of conductivity spectra determined by equation 6.7 (Di Noto et al., 2002b): σ ′(ω ) =
Z′ {k ⋅[Z ′(ω )]2 + [Z ′′(ω )]2 }
[6.7]
Z is the complex impedance of the system and k = l/S is the cell constant; l and S are the length and the cross-section of the ion conductor. This relationship is easily determined by considering that: σ* (ω ) =
1 1 = ρ* (ω ) [ k ⋅ Z* (ω )]
[6.8]
where σ*(ω), ρ*(ω) and Z*(ω) are the complex conductivity, the complex resistivity and the complex impedance, respectively. The experimental profile of σ′(ω) vs. ω can be interpolated by using the UPL phenomenological equation 6.2, provided that it shows a single relaxation event. The UPL equation 6.2 can be interpreted on the basis of the jump relaxation model. This model was originally developed to describe inorganic solid ionic conductors such as RbAg4I5, Na-β-alumina and AgBr below the limit of 100 GHz (Funke et al., 1992; Cramer et al., 1995). The UPL equation cannot be simply explained by hypothesizing a random hopping of the charge carriers in a lattice of mostly empty sites. Random hopping would imply that the charge carriers do not affect one another and it would be possible to write equation 6.9 for a single mobile species: ⎧⎪ σ* (ω ) ∝ FT ⎨ ⎩⎪
N
∑ ν ( 0 ) ⋅ ν (t ) i
i, j
j
⎫⎪ ⎬ ≈ FT {N ⋅ δ (t )} ⎪ hops ⎭
[6.9]
In this case, the complex conductivity σ*(ω) is proportional to the Fourier transform (FT) of the correlation function of the velocity vectors of the
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Hybrid inorganic–organic polymer electrolytes
267
In[σ´(ω)] In σ´(∞)
In σ´(0) ω ω2 = t2–1
ω1 = t1–1
6.49 Typical profile of the real part of the complex conductivity vs. ω in solid electrolytes characterized by a disordered structure.
carriers involved in inter-site hopping (hops) (Funke, 1991). Equation 6.9 predicts a linear correlation between ln[(σ′(ω)] and ω, which is not observed. The jump relaxation model includes the possibility of an effective correlation and memory of the charge carriers; then, it is possible to write equation 6.10: ⎧⎪ σ* (ω ) ∝ FT ⎨ ⎩⎪
N
∑ ν ( 0 ) ⋅ ν (t ) i
j
i, j
⎫⎪ ⎬ ≈ FT νi ( 0 ) ⋅ ν j (t ) ⎪ hops ⎭
{
hops
}
[6.10]
Equation 6.10 predicts correctly that the profile of ln[(σ′(ω)] vs. ω is sigmoidal, as shown in Fig. 6.49. The potential V(x) of the ionic carrier is the sum of the periodic lattice potential Vp(x), which is sinusoidal, and the coulombic potential VC(x), which is approximated as parabolic as shown by Funke and Riess (1984) (see Fig. 6.50a): V ( x ) = Vp( x ) + VC ( x ) z2 ⋅ e 2 1⎞ ⎛ 1 VC ( x ) = − ⎟ ∑ ⎜ 4π ⋅ ε j ⎝ rj − r rj ⎠
= VC*( x ) ⋅ r 2 + L ( j = 1, . . , N ions ) configurat ions
[6.11] z · e is the charge of the carriers and L indicates higher order terms which can be ignored (here and hereafter). From time to time, the charge carrier in the filled state A is thermally activated and hops into a vacant neighbouring site B. The charge carrier may now either hop back relatively soon (this occurrence is energetically possible) or stay at its new position. If the charge carrier settles in, the medium reorganizes and the coulomb potential will move towards site B, turning the empty site B into a filled state A. These events are summarized on Fig. 6.50(b). If the charge carrier is in A at t = 0, the probability to find it in an empty site B at the time t is: WB(t) = W0(t) + W2(t) + L
© Woodhead Publishing Limited, 2010
[6.12]
268
Polymer electrolytes (a) Vp(x)
VC(x)
V(x) δ B
Δ
A (b)
δ Δ
t1
δ(t)B→C
δ(t)B→A
x(t)
x0 A
x0 B
C
6.50 Components of the overall potential (a) and hopping dynamics (b) in the jump relaxation model.
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Hybrid inorganic–organic polymer electrolytes
269
and the probability to find it in A is WA(t) = W1(t) + L
[6.13]
W0(t) is the probability that the charge carrier is still in B at the time t, W1(t) is the probability of finding the charge carrier in A at the time t after a back-hopping from B and W2(t) is the probability of finding the charge carrier in B after a forward-hopping followed by a back-hopping. On the basis of the linear response theory of Kubo and the theorem of fluctuationdissipation it is possible to write equation 6.14 (Kubo, 1957; Funke and Riess, 1984): ∞
σ ′(ω ) = σ ′( ∞ ) ⋅ ∫ {δ ( t ) + WB′(t )} ⋅ cos (ωt ) dt
[6.14]
0
The term WB(t) can be calculated directly using the probability theory or can be obtained by the correlation function of the velocity vectors of the ions. If the cross terms are neglected as in the approximation underlying the Nernst–Einstein (NE) equation, it is possible to write equation 6.15 (Funke, 1991):
(1 N ) ⋅
N
∑ νi ( 0 ) ⋅ ν j ( t ) i, j
≈ ν ( 0 ) ν (t )
NE hop
=
hop
1 d2 2 ⋅ r (t ) 2 dt 2
hop
[6.15]
The mean squared displacement 〈r2(t)〉hop of each charge carrier can be related to WB(t) by the following equation 6.16: r 2(t )
t
hop
x0 2
= Γ 0 ⋅ ∫ WB(t ′ ) dt ′
[6.16]
0
where x0 is the A–B distance between two neighbouring sites and (Funke and Riess, 1984): Γ0 = ν0C · exp[−(δ/kBT)]
[6.17]
is the back-hopping rate through the energy barrier δ with the attempt frequency ν0C. Therefore equation 6.18 is obtained (Funke, 1987) where W′B(t) indicates the first derivative of WB(t): v ( 0 ) ⋅ v (t )
hop
=
x0 2 Γ 0{δ ( t ) + WB′( t )} 2
[6.18]
By substituting 6.18 in 6.10, it is found that: σ ′(ω ) = N
∞
x02 Γ 0 ⋅ ∫ {δ (t ) + WB′(t )} ⋅ cos (ωt ) dt 2 0
[6.19]
where σ′(∞) is equal to the factor outside of the integral. A successful jump forward of a charge carrier from a site A to a site B requires a site relaxation
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Polymer electrolytes
following the hopping event before the back-hopping event can occur. The initial back-hopping rate is defined as 1/τ* while the initial site-relaxation rate is 1/τ2. The times τ2 and τ1 are the reciprocal of the frequency points read on the sigmoid ln[(σ′(ω)] vs. ω before and after the flexus point, respectively (see Fig. 6.49). If we indicate with WA(t) the probability of backhopping it is possible to write (Cramer et al., 1995):
(
dWA (t ) 1 =− τ* dt
)
[6.20]
t =0
In addition,
(
1 1 dδ back(t ) = τ 2 kBT dt
)
[6.21]
t =0
where δback(t) is the energy barrier preventing back-hopping. It is demonstrated that, in the UPL equation 6.2, p = τ2/τ*. For small values of t the site-relaxation is not complete and it is possible to write: δ (t ) −WA′ (t ) = WA ( t ) ⋅ ν0 B ⋅ exp ⎡⎢ − back ⎤⎥ ⎣ kBT ⎦
[6.22]
In addition, −
(δ (t ) − δ ) ⎤ WA′ ( t ) 1 = exp ⎡⎢ − back ⎥⎦ WA ( t ) τ* kBT ⎣
[6.23]
where −δ ⎤ 1 = ν0 B exp ⎡⎢ − τ* ⎣ kBT ⎥⎦
[6.24]
The initial equation 6.23 for 1/τ* is obtained as WA(0) = 1 and δback(t) = δ. The rate of back-hopping is regulated by the temporal dependence of δback(t). If there is no site-relaxation, δback(t) = δ for every t and the integration of equation 6.23 yields equation 6.25: t WA ( t ) = exp ⎡ − ⎤ ⎢⎣ τ* ⎥⎦
[6.25]
However, numerical simulations have shown that for short values of t (high frequencies) the rate of site-relaxation is well described by the following relationship (Maass et al., 1991): ⎡ − −δ back(t ) ⎤ = 1 ⎢⎣ kBT ⎥⎦ τ 2 + t
[6.26]
This equation coincides with equation 6.21 as t → 0. The integration of equation 6.26 yields:
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Hybrid inorganic–organic polymer electrolytes t δ back(t ) − δ = kBT ⋅ ln ⎡⎢1 + ⎤⎥ ⎣ τ2 ⎦
271 [6.27]
As equation 6.27 is substituted in equation 6.23, the following result is obtained: ⎡ WA′ ( t ) 1 ⎢ 1 − = ⎢ WA ( t ) τ* ⎢ 1 + t τ2 ⎣
⎤ ⎥ ⎥ ⎥ ⎦
[6.28]
The integration of equation 6.28 yields: t WA ( t ) = ⎛ 1 + ⎞ ⎝ τ2 ⎠
τ − 2 τ*
[6.29]
If τ2 → ∞, the original exponential form is obtained: t WA ( t ) = lim ⎛ 1 + ⎞ τ 2 →∞ ⎝ τ2 ⎠
τ − 2 τ*
t = exp ⎡ − ⎤ ⎢⎣ τ* ⎥⎦
[6.30]
By using the formula of Kubo: ∞
⎡ ⎤ σ ′(ω ) − σ ′( 0 ) = [σ ′( ∞ ) − σ ′( 0 )]⋅ ⎢1 + ∫ WA ( t ) ⋅ cos (ωt ) dt ⎥ ⎣ 0 ⎦
[6.31]
and using equation 6.30 it is found that, at high frequencies, equation 6.32 holds true: −p
σ ′(ω ) − σ ′( 0) = [ σ ′( ∞) − σ ′( 0)]⋅ [1 + (ωτ 2 )−1 ]
[6.32]
with p = τ2/τ*. The value of p can vary between 0 and 2, as required by the transformation conditions. This equation is satisfactory for the simulation of the plateau at high frequencies. At the lower frequencies (Fig. 6.49), it is necessary to consider the case ωτ2 << 1; in that instance, equation 6.32 can be expressed as reported below (equation 6.33): σ ′(ω ) = σ ′( 0 ) +
σ ′( 0 ) ⋅[ σ ′( ∞ ) − σ ′( 0 )] p ⋅ (ωτ 2 ) σ ′( 0 )
[6.33]
If 1
⎡ σ ′( ∞ ) − σ ′( 0 ) ⎤ p τ1 = ⎢ ⎥⎦ ⋅ τ 2 σ ′( 0 ) ⎣
[6.34]
then equation 6.33 becomes the UPL equation 6.2. This latter condition is plausible given the geometrical meaning of τ2 and τ1 as obtained from the plot of the sigmoid ln[σ′(ω)] vs. ω. The τ2 and τ1 times are the limits of the
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region connecting the σ′(0) and σ′(∞) plateaus and can be both considered site-relaxation times. τ1 can be directly obtained from the fitting of the data with the UPL equation. The analysis of the real part of the conductivity spectra also allows one to gain information on the conductivity mechanism of the material. By carrying out measurements at different temperatures it is possible to plot σ′, determined on the low frequency plateau of ln[σ′(ω)], vs. 1/T. The resulting data can be fitted with suitable functions (e.g. Arrhenius, VTF . . .), which yield the activation energy of the conductivity for the material. The same procedure can be carried out to evaluate the activation energy of τi, i = 1, 2; the latter values are determined starting from the analysis of ln[σ′(ω)] profiles as shown on Fig. 6.49. If the activation energies of σ′ and τi are approximately the same, it can be concluded that the ratedetermining step of the conductivity mechanism (expressed by σ′) is the hopping process of the charge carriers (described by τi). On the other hand, if the activation energy of τi is much smaller with respect to σ′, it can be concluded that even if the charge carriers are free to hop quickly, the overall long-range charge transport is slowed by some other mechanism. Finally, if the activation energy of τi is much larger with respect to σ′, then the overall charge transport mechanism is not regulated by ion hopping but by some other processes, e.g. the segmental motion of the polymer chains of a macromolecular material.
6.6
Conclusions
The various examples of hybrid inorganic–organic polymer electrolytes presented in this chapter clearly show that there is an enormous number of feasible alternatives in the quest to devise good materials for application as ionic conductors. In particular, the inorganic species proved to be very flexible networking centres and provide good control of the nanostructure of the hybrid materials. The interactions between the inorganic nanofillers and the host polymer lead to a significant modulation of the thermal, mechanical, structural and electric properties of the resulting nanocomposite systems. A highly detailed characterization of all the proposed materials under the chemical, morphological, structural and electrical points of view is absolutely crucial to obtain a clear picture of the large number of factors controlling the performance of hybrid inorganic–organic PEs. In particular, extensive studies by electrochemical impedance spectroscopy (EIS) and broadband dielectric spectroscopy (BDS), followed by significant efforts in the interpretation of the resulting data are essential tools to reach this target. Indeed, the examples presented in this chapter indicate that it is possible to devise materials showing the desired properties and characterized by high chemical, thermal and electrochemical stability and good ionic conductivity with a variety of charge carriers (e.g. lithium and magnesium
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ions, protons). Further studies will be necessary to gain a greater understanding of the structure–performance relationship in these hybrid inorganic–organic materials. The latter are among the most promising candidates for future application as highly efficient ionic conductors in a variety of advanced devices including energy storage and conversion systems such as high energy density batteries and PEMFCs.
6.7
Acknowledgements
The authors would like to extend their most sincere thanks to Julianne Vittadello for her precious contribution during the reviewing process of the chapter.
6.8
References
abraham k m, alamgir m (1990), ‘Li+-conductive solid polymer electrolytes with liquid-like conductivity’, J Electrochem Soc, 137, 1657–1658. angell c a (1997), ‘Entropy and fragility in supercooling liquids’, J Res Natl Inst Stand Technol, 102, 171–185. appleby a j (1999), ‘The electrochemical engine for vehicles’, Sci Am, 281, 74–79. armand m, gorecki w, beranger s, michot c (2002), ‘Polymer electrolytes: ancestors destined to industrial oblivion & fundamental studies or a renaissance?’, Book of Abstracts of the 8th International Symposium on Polymer Electrolytes, Santa Fe, New Mexico, 19–24 May, 1. biscazzo s, vittadello m, lavina s, di noto v (2002), ‘Synthesis and structure of electrolytic complexes based on α-hydro-ω-oligo(oxyethylene)hydroxypoly[oligo(oxyethylene)oxydimethylsililene] and δ-MgCl2’, Solid State Ionics, 147, 377–382. blonsky p m, shriver d f, austin p, allcock h r (1986), ‘Complex formation and ionic conductivity of polyphosphazene solid electrolytes’, Solid State Ionics, 18, 258–264. bouridah a, dalard f, deroo d, cheradame h, le nest j f (1985), ‘Poly(dimethlsiloxane)–poly(ethylene oxide) based polyurethane networks used as electrolytes in lithium electrochemical solid state batteries’, Solid State Ionics, 15, 233–240. chadwick b m, jones d w, wilde h j, yerkess j (1985), ‘X-ray and neutron diffraction studies of the crystal structures of the dicesium lithium hexacyanometallates of iron(III) and cobalt(III)’, J Cryst Spectrosc, 15, 133–146. chung s h, wang y, greenbaum s g, golodnitsky d, peled e (1999), ‘Uniaxial stress effects in poly(ethylene oxide)-LiI polymer electrolyte film A 7Li nuclear magnetic resonance study’, Electrochem Solid State Lett, 2, 553–555. cleghorn s j c, ren x, springer t e, wilson m s, zawodzinski c, zawodzinski t a, gottesfeld s (1997), ‘PEM fuel cells for transportation and stationary power generation applications’, Int J Hydrogen Energ, 22, 1137–1144. cotton f a, murillo c a, bochmann m (1999), Advanced Inorganic Chemistry, 6th Edition, New York, Wiley Interscience. cramer c, funke h, saatkamp t, wilmer d, ingram m d (1995), ‘High-frequency conductivity plateau and ionic hopping processes in a ternary lithium borate glass’, Z Naturforsch, 50, 613–623.
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croce f, brown s d, greenbaum s g, lane s m, salomon m (1993), ‘Lithium-7 NMR and ionic conductivity studies of gel electrolytes based on poly(acrylonitrile)’, Chem Mater, 5, 1268–1272. dias f b, plomp l, veldhuis j b j (2000), ‘Trends in polymer electrolytes for secondary lithium batteries’, J Power Sources, 88, 169–191. di noto v (1997), ‘A novel polymer electrolyte based on oligo(ethylene glycol) 600, K2PdCl4 and K3Fe(CN)6’, J Mater Res, 12, 3393–3403. di noto v (2000), ‘Zeolitic inorganic–organic polymer electrolyte based on oligo(ethylene glycol) 600 K2PdCl4 and K3Co(CN)6’, J Phys Chem B, 104, 10116–10125. di noto v (2002), ‘Electrical spectroscopy studies of lithium and magnesium polymer electrolytes based on PEG400’, J Phys Chem B, 106, 11139–11154. di noto v, vittadello m (2005), ‘Two new siloxanic proton-conducting membranes Part I. Synthesis and structural characterization’, Electrochim Acta, 50, 3998–4006. di noto v, zago v (2004), ‘Inorganic–organic polymer electrolytes based on PEG400 and Al[OCH(CH3)2]3 I. Synthesis and vibrational characterizations’, J Electrochem Soc, 151, A216–A223. di noto v, furlani m, lavina s (1996a), ‘Synthesis, characterization and ionic conductivity of poly[(oligoethylene oxide) ethoxysilane] and poly[(oligoehylene oxide) ethoxysilane]/(EuCl3)0.67’, Polym Adv Technol, 7, 759–767. di noto v, bettinelli m, furlani m, lavina s, vidali m (1996b), ‘Conductivity, luminescence and vibrational studies of the poly(ethylene glicol) 400 electrolyte based on europium trichloride’, Macromol Chem Phys, 197, 375–388. di noto v, barreca d, furlan c, armelao l (2000), ‘Zeolitic inorganic–organic polymer electrolytes: a material based on poly(ethylene glycol) 600, SnCl4 and K4Fe(CN)6’, Polym Adv Technol, 11, 108–121. di noto v, fauri m, vittadello m, lavina s, biscazzo s (2001), ‘Zeolitic inorganic– organic polymer electrolytes: synthesis, characterization and ionic conductivity of a material based on oligo(ethylene glycol) 600, (CH3)2SnCl2 and K4Fe(CN)6’, Electrochim Acta, 46, 1587–1594. di noto v, vittadello m, pace g, biscazzo s, lavina s (2002a), ‘Synthesis and characterization of [PEG400-alt-DEOS]/(δ-MgCl2)0.2597 complex’, Macromol Chem Phys, 203, 1201–1210. di noto v, fauri m, vittadello m, lavina s, biscazzo s (2002b), ‘Conductivity, thermal stability and morphology of a new Z-IOPE inorganic–organic network with the formula [FexSny(CH3)2y(CN)zClv(C2nH4n+2On+1)Kl]’, Macromol Chem Phys, 203, 354–362. di noto v, zago v, biscazzo s, vittadello m (2003a), ‘Hybrid inorganic–organic polymer electrolytes: synthesis, FT-Raman studies and conductivity of {Zr[(CH2CH2O)8.7]ρ/(LiClO4)z}n network complexes’, Electrochim Acta, 48, 541–554. di noto v, vittadello m, lavina s, biscazzo s, fauri m (2003b), ‘The first lithium zeolitic inorganic–organic polymer electrolyte based on PEG600, Li2PdCl4 and Li3Fe(CN)6: part I, synthesis and vibrational studies’, Electrochim Acta, 48, 2047–2058. di noto v, zago v, pace g, fauri m (2004a), ‘Inorganic–organic polymer electrolytes based on PEG400 and Al[OCH(CH3)2]3 II. Morphology, thermal stability and conductivity mechanism’, J Electrochem Soc, 151, A224–A231.
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di noto v, vittadello m, greenbaum s g, suarez s, kano k, furukawa t (2004b), ‘A new class of lithium hybrid gel electrolyte systems’, J Phys Chem B, 108, 18832–18844. di noto v, vittadello m, jayakody j r p, khalfan a n, greenbaum s g (2005), ‘Two new siloxanic proton conducting membranes Part II. Proton conductivity mechanism and NMR study’, Electrochim Acta, 50, 4007–4014. di noto v, gliubizzi r, negro e, pace g (2006), ‘Effect of SiO2 on relaxation phenomena and mechanism of ion conductivity of [Nafion/(SiO2)x] composite membranes’, J Phys Chem B, 110, 24972–24986. di noto v, gliubizzi r, negro e, vittadello m, pace g (2007), ‘Hybrid inorganic– organic proton conducting membranes based on Nafion and 5 wt. % of MxOy (M = Ti, Zr, Hf, Ta and W) Part I. Synthesis, properties and vibrational studies’, Electrochim Acta, 53, 1618–1627. di noto v, piga m, piga l, polizzi s, negro e (2008), ‘New inorganic.organic proton conducting membranes based on Nafion® and [(ZrO2)·(SiO2)0.67] nanoparticles: synthesis, vibrational studies and conductivity’, J. Power Sources, 178, 561–574. di noto v, lavina s, negro e, vittadello m, conti f, piga m, pace g (2009), ‘Hybrid inorganic–organic proton conducting membranes based on Nafion and 5 wt% of MxOy (M = Ti, Zr, Hf, Ta, and W). Part II: Relaxation phenomena and conductivity mechanism’, J. Power Sources, 187, 57–66. di noto v, negro e, lavina s (2010a), ‘Broadband dielectric spectroscopy and conductivity mechanism of Nafion 117 and Nafion/[ZrO2] hybrid inorganic–organic membranes’, in A. Herring, Fuel Cell Membranes; ACS Symposium Series, American Chemical Society, Washington, DC. di noto v, piga m, lavina s, negro e, yoshida k, ito r, furukawa t (2010b), ‘Structure, properties and proton conductivity of Nafion/[(TiO2)·(WO3)0.148]ψTiO2 nanocomposite membranes’, Electrochim Acta, 55, 1431–1444. feuillade g, perche p (1975), ‘Ion-conductive macromolecular gels and membranes for solid lithium cells’, J Appl Electrochem, 5, 63–69. funke k (1987), ‘Debye–Hückel-type relaxation processes in solid ionic conductors Part II: Experimental evidence’, Z Phys Chem Neue Fol, 154, 251–295. funke k (1991), ‘Is there a “universal” explanation for the “universal” dynamic response?’, Ber Bunsen Phys Chem, 95, 955–964. funke k, riess i (1984), ‘Debye-Hückel-type relaxation processes in solid ionic conductors Part I: The model’, Z Phys Chem Neue Fol, 140, 217–232. funke f, kloidt t, wilmer d (1992), ‘Jump relaxation in RbAg4I5 by dynamic conductivity and quasielastic neutron scattering’, Solid State Ionics, 53–56, 947–954. furukawa t, imura m, yuruzume h (1997), ‘Broad-band conductive spectra of polypropylene oxide complexed with LiClO4’, Jpn J Appl Phys 1, 36, 1119–1125. gadjourova z, andreev y g, tunstall d p, bruce p g (2001), ‘Ionic conductivity in crystalline polymer electrolytes’, Nature, 412, 520–523. gray f m (1997), Polymer Electrolytes, RSC Materials Monographs, Cambridge, Royal Society of Chemistry. guyomard d, tarascon j m (1994), ‘Rocking-chair or lithium-ion rechargeable lithium batteries’, Adv Mater, 6, 408–412. huq r, koksbang r, tonder p e, farrington g c (1992), ‘Effect of plasticizers on new ambient temperature polymer electrolyte’, Electrochim Acta, 37, 1681–1684.
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theisen t (2003), ‘Natural gas for power generation and the automotive market,’ in Vielstich W, Lamm A, Gasteiger H A, Handbook of Fuel Cells: Fundamentals, Technology and Applications, Chichester, John Wiley & Sons, vol. 3, 25–38. vittadello m, biscazzo s, lavina s, fauri m, di noto v (2002), ‘Vibrational studies of the ion-polymer interactions in α-hydro-ω-oligo(oxyethylene) hydroxypoly[oligo(oxyethylene)oxydimethylsililene]/δ-MgCl2’, Solid State Ionics, 147, 341–347. vittadello m, suarez s, chung s h, fujimoto k, di noto v, greenbaum s g, furukawa t (2003), ‘The first lithium zeolitic inorganic–organic polymer electrolyte based on PEG600, Li2PdCl4 and Li3Fe(CN)6: Part II, thermal stability, morphology and ion conduction mechanism’, Electrochim Acta, 48, 2227–2237. vittadello m, negro e, lavina s, pace g, safari a, di noto v (2008), ‘Vibrational studies and properties of hybrid inorganic–organic proton conducting membranes based on Nafion and hafnium oxide nanoparticles’, J Phys Chem B, 112, 16590–16600. wright p v (1998), ‘Polymer electrolytes – the early days’, Electrochim Acta, 43, 1137–1143. wright p v, zheng y, bhatt d, richardson t, ungar g (1998), ‘Supramolecular order in new polymer electrolytes’, Polym Int, 47, 34–42.
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7 Using nuclear magnetic resonance spectroscopy in polymer electrolyte research S. A B B R E N T, University of South Bohemia, Czech Republic, and S. GREENBAUM, Hunter College of the City University of New York, USA
Abstract: The versatility and power of NMR spectroscopy to probe matter at the atomic level and the possibility of achieving detailed knowledge and understanding about the dynamic processes governing ionic and segmental motion within polymer electrolyte materials will be described. Many examples of results reported from different scientific groups active in the field of polymer electrolyte research will be presented. Key words: NMR techniques, polymer electrolytes.
7.1
Introduction
Nuclear magnetic resonance (NMR) spectroscopy measurements have been firmly established among techniques used for obtaining valuable information about chemical and physical properties in a wide variety of scientific fields from biology and medicine to chemistry and physics. This technique is element-specific, dependent on single nuclear species having a magnetic moment. The sample being examined is placed inside a strong, usually superconducting magnet and according to the aim of the study the sample is exposed to sequences of radiofrequency pulses. Depending on the atomic surroundings, the response from the magnetized nuclei will be very individual, even unique. Analysis of the number, shape, size and position of bands in the resulting spectrum reveal information about chemical and physical properties of the sample.1–3 An important feature of NMR is that the integrated spectroscopic signal is proportional to the number of resonant NMR nuclei in the sample. Thus, NMR analysis can be used to ascertain the relative phase content of a material in cases where individual spectral contributions can be resolved. The NMR-active nuclei can be divided into two types, the spin-1/2 nuclei such as 13C and 1H and the quadrupolar nuclei (spin > 1/2) such as 6Li, 7Li, 59Co, and 27Al. Each local environment in a solid gives rise to its own distinct resonance, and as long as 278 © Woodhead Publishing Limited, 2010
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the environments are sufficiently different, these local environments can be resolved and quantified. The resolution of the spectrum will depend on the nature of the interactions that control, shift, and broaden the resonances. In a diamagnetic material, the nature of the coordination environment for lithium (6Li or 7Li) will give rise to a distinct and characteristic ‘chemical shift.’ Unfortunately, Li is a light element, and as in the case of 1H NMR, its chemical shift range is small, which can result in overlapping resonances. NMR spectroscopy may be used to probe processes with a wide range of correlation times. Spin–lattice relaxation (T1) measurements are sensitive to motional processes occurring at time scales of the reciprocal of the NMR frequency (several to tens of nanoseconds), and other relaxation times (T1ρ and T2) and sample millisecond scale correlation times. Similarly, the observation of spectral line-narrowing signifies motion comparable to the reciprocal of the line width (i.e. milliseconds for kilohertz linewidths). In addition, self-diffusion coefficients (D) of mobile species (D > 10−9 cm2/s) can be determined by pulsed gradient spin-echo (PGSE) methods, as originally described by Stejskal and Tanner, to be described in detail later. This is an important tool in studies of ion transport in gel and polymer electrolytes, in that both cation (i.e. 7Li) and anion (for anions containing 19F) diffusion coefficients can be obtained.
7.1.1 Polymer electrolytes The need for an efficient, compact and environmentally friendly energy storage source has led to the evolution of all solid battery systems that utilize electrolytes consisting of solid or at least non-leaking polymer materials. Properties of these are often compromises achieved by optimization of mechanical, dynamic and compatibility issues with addition of economic and environmental impact aspects. There have been many different approaches to this problem described in the literature,4–8 where the research moved from solid polymer (i.e. high molecular weight or cross linked electrolytes9 to polymer matrix hosting liquid electrolytes, gel electrolytes10 and later composites containing filler additives.11–16 In the past years the use of nanoscopic additives in polymer electrolytes has brought many interesting results, which will be described in more detail later.17–24
7.2
Nuclei possibility
Essentially every element in the periodic table has at least one isotope with a nuclear magnetic moment. Not all can be easily observed by the NMR
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technique as the receptivity of different nuclei varies greatly, depending on both the nucleus’s gyromagnetic ratio, its natural abundance, and other factors such as quadrupole moment. The frequency at which a response from a certain nucleus can be observed is also dependent on the strength of the magnetic field. With optimized equipment and experimental settings a wide range of different species can be studied. Furthermore, it is often possible to purchase chemicals enriched with isotopes that do not occur naturally in high enough abundance, common examples being 2H and 13C. Most easily measured is the highly abundant, sensitive, and stable nucleus of hydrogen 1H, followed by only slightly less receptive fluorine 19F. Deuterium 2D is used as an alternative to proton measurements, and deuterated samples are often used as standards in organic analysis. Considering the importance of carbon in organic and polymer chemistry, one of the most often used nuclei is 13C. It is much less sensitive but has other advantages compared with the proton, such as a much larger chemical shift range and, because of the much lower abundance as compared to the proton (1.1%), practically no observed spin–spin couplings, interactions between nuclei of the same type. The resulting spectrum is therefore greatly simplified and much less crowded than a proton one, yielding sharp signals in a wide chemical shift range, which means that the individual bands are easily identified. Lithium, often used as cation in polymer electrolytes, has two useful nuclei. 6Li yields sharp signals but has low sensitivity due to its low natural abundance (6%) and lower (than 7Li) gyromagnetic ratio. On the other hand, 7Li is highly sensitive but has a larger quadrupole moment so its signals are broader. Both nuclei exhibit only a modest chemical shift range. Other isotopes often studied in polymer electrolyte research are those of different metal cations. In the literature, measurements on sodium 23Na, magnesium 25Mg, potassium 39,40K, copper 63,65Cu, nickel 61Ni, aluminum 27Al and many others can be found. Besides fluorine many other measurable species occur in polymer electrolyte materials as anions, for example 31P, 11B or 35,37Cl. Finally, important nuclei for biochemical research, but also to lesser extent for this field are 17O and 14,15N. Owing to this richness of measurable nuclei, it is often possible to examine several species in one single sample, and usually, if choosing carefully, many if not all the different components of the polymer electrolyte mixture can be observed separately and mutually compared.25,26 What follows is a nonexhaustive review of previous NMR investigations of polymer electrolytes representing a broad range of approaches and materials systems. This chapter is restricted to lithium-containing polymers and does not discuss proton-conducting polymer electrolytes for fuel cell applications.
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Liquid state nuclear magnetic resonance
5.0
4.5
4.0
3.5
3.0
2.5
1.030
1.341 1.234
2.089
3.358 3.099 2.396 2.736 2.607
3.641
4.484
In liquid samples, because of the interactions that are averaged to zero by the rapid tumbling of individual molecules on the picosecond timescale, the resulting spectra show sharp, narrow bands. Their position, the so-called chemical shift, is determined by the local magnetic field arising from the chemical bonds in the vicinity of the nuclei and their interaction with the laboratory field. The resulting spectra can therefore be used for identification of unknown chemical species, or for clarifying unknown chemical structures. This was performed on a new class of difluoroalkoxyborane compounds later used as additives in a polymer electrolyte material.26 Zheng et al.27 used 1H NMR to solve molecular structure of a newly synthesized co-polymer. Wang and Tang28 compared 1H NMR spectra of a polymer in solution (Fig. 7.1) with Fourier transform infrared (FTIR) spectroscopic results, when characterizing novel gel polymer electrolyte material based on a blend of several polymeric species. Besides the use of novel materials reported by many research groups, polymer electrolytes consisting of even well-known chemical structures present significant scientific challenges because interactions between the components are often not entirely clear. It is these interactions that facilitate or impede the desired ionic mobility within the materials, and knowledge and understanding of these is therefore of major importance. The great sensitvity of high resolution NMR to the exact surrounding of the observed species makes it possible to explore molecular interactions in a unique way. As an example, the interactions that exist between added plasticizer and a new polymer matrix, examined by high resolution NMR, can be elucidated29 (see Fig. 7.2). Polymer electrolytes are solid or semi-solid systems not easily examined properly at high resolution. For the purpose of revealing interactions
2.0
1.5
1.0
0.5
ppm
7.1 1H NMR spectrum of PMAML. From: Z-l. Wang, Z-Y. Tang, Electrochimica Acta 2004, 49, 1063–1068.28
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(c)
(b)
(a)
200 180 160 140 120 100 80 d (ppm)
60
40
20
0
7.2 13C -NMR spectra of (a) PC (liquid), (b) VA–PSVE–Na (solid), (c) PC/ VA–PSVE–Na = 1.5/3.0 (gel). From: L-Y. Tian, X-B. Huang, X-Z. Tang, European Polymer Journal 2004, 40, 735–42.29
between individual species, model systems are often used. The classical long-chained polymers are replaced by polymers with short chains or even oligomers and liquid electrolytes are obtained. High molecular weight poly(ethylene oxide) (PEO) has often been replaced by so-called glymes, basically short chain PEO molecules,30–32 as they behave similarly to amorphous-phase PEO when end-capped with alkyl groups, shown both experimentally33–35 and by theoretical calculations.36–38 Ion-pair formation constants for different salts with fluorine-containing anions have been estimated by 19 F NMR for several glymes and successfully compared to data obtained from conductivity measurements.39
7.4
Solid state nuclear magnetic resonance
To study condensed and relatively rigid systems, solid state NMR cannot give as clear a picture of the composition of an examined sample due to resolution limitations compared to the liquid state. Nevertheless, magic angle spinning (MAS) methods do succeed in offering moderately high resolution and this will be discussed more fully later. Although the identification of individual species is more limited, the dynamics of the system as a whole can be probed from line width, line shape and relaxation behavior of the resulting peaks as described by Bloembergen et al.,40 and many others since that seminal work. As mobility within an electrolyte is of major
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importance, solid state NMR experiments are of great use for the assessment of dynamic properties of these materials.
7.4.1 Morphology Polymer electrolytes are heterogeneous materials with complicated phase diagrams usually consisting of both crystalline and amorphous components.41,42 Spectra of pure PEO and dilute salt PEO complexes typically show a broad and a narrow NMR signal that are superimposed. For pure PEO, these components have been assigned to the amorphous and the crystalline domains, but the assignments have varied. Connor and Hartland43 originally used proton spin–lattice relaxation to study the heterogeneous nature of pure PEO. The broad component, with a fast relaxation rate, was assigned to the amorphous and the narrow component, with longer relaxation rates, to the crystalline component. This assignment was later reversed by Dechter44 who assigned the broad component to the crystalline and the narrow component to the amorphous based on a reanalysis of the proton and 13C rotating frame spin–lattice relaxation data (T1ρ, see below). Johansson and Tegenfeldt45 reanalyzed and agreed with the results found by Dechter. All NMR studies of PEO electrolytes since have used this assignment. However, pure PEO and PEO electrolytes are fundamentally different samples and NMR assignments of their bands are not always straightforward.46 High salt concentrations have been shown to retain only the narrow component, while the more dilute PEO–salt samples contain both the narrow and broad components.47–49 In dilute PEO:salt electrolytes, three regions have been shown to exist: (1) pure crystalline PEO, (2) amorphous PEO, (3) amorphous PEO–Li salt complex, and (4) crystalline PEO–Li salt complex.42,50
7.4.2 Thermal properties Variable temperature study with a subsequent line shape analysis enables a clear insight of the temperature-dependent dynamics of the system, where the glass transition temperature (Tg) of the amorphous phase, melting point or a crystallization onset of a present crystalline phase, or any other changes caused by altered dynamics of the involved species can be assessed. This was recognized by early investigations of polymer electrolytes and will be discussed in greater detail later. Linewidth changes dependent on salt amount were examined by Liao et al.51 (see Fig. 7.3), and the line shape was also examined on heating and cooling, as shown in Fig. 7.4. From these graphs it is possible to conclude that spectral lines are narrow in the liquid crystal phase, while they broaden
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7.3 7Li and 19F NMR linewidths versus temperature for C18O5: LiBF4 (1 : x) second heating cycle data, Tm denotes the side chain melting temperature. Filled symbols are for 7Li and open symbols are for 19F. (a) 䉬, 䉫, x = 0.7, ‘salt-deficient’ sample; (b) 䊏, x = 1, ‘equimolar’ sample; (c) 䉱, 䉭, x = 1.3 ‘excess salt’ sample; (d) selected 7Li NMR signals from (b). From: Y-P. Liao, D. C. Apperley, J. Liu, Y. Zheng, P. V. Wright, Electrochimica Acta 2007, 53, 1444–54.51
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7.4 (a) C18O1 : LiBF4 (1 : 0.6); 䉱, 䉭, 7Li and 19F NMR linewidths for the second heating and cooling cycles. Filled symbols are for 7Li and open symbols are for 19F. (b) C16O1O5(21%) : LiBF4 (1 : 0.7); 䊏, 䊐, 7Li and 19F NMR linewidths of the signals in C16O1O5(21%) : LiBF4 (1 : 0.7) for the second heating and cooling cycles. (c) 7Li lineshape at −7 °C; (d) 7Li lineshape at −20 °C; (e) 19F line shape at −70 °C. Tm denotes the side chain melting temperature. (f) 19F line shape at −20 °C. From: Y-P. Liao, D. C. Apperley, J. Liu, Y. Zheng, P.V. Wright, Electrochimica Acta 2007, 53, 1444–54.51
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considerably as the material crystallizes, indicating a decrease in dynamics, in accordance with both thermal analysis and conductivity measurements. The influence of salt concentration is also easily observed in the curvature of the graph lines, where the onset of increase of the linewidths moves to lower temperatures, indicating higher dynamics of the material at low temperatures. The difference between 7Li and 19F spectral linewidths is partly due to their different gyromagnetic ratios, but also indicates differences between the ionic species’ mobilities. Often, sets of samples are studied where one parameter, concentration or relative amount of one of the components, or one component itself, is varied. The observed changes in dynamics and morphology caused by the altered composition of the sample, such as relative amount of crystallinity, altered coordination of the involved species, or different degree of ionpairing, are of paramount importance in the quest for an optimized electrolyte system.
7.4.3 Composition The complexity of polymer electrolyte systems and the influence of addition of different nanoparticles were discussed by Singh and Bhat52 based on 7Li linewidth observation. They state that while it is generally expected that when ionic motion occurs, the NMR linewidth of the mobile nucleus decreases, several discrepancies have been observed. For example, Dai et al.17 with (PEO)1.5 + LiI + nanoscale Al2O3 or MgO, Chung et al.21 with (PEO)8 + LiClO4 + TiO2 nanoparticles and Forsyth et al.18 with 3PEG + LiClO4 + TiO2, SiO2 and Al2O3 nanoparticles observed an increase in the 7 Li linewidth on addition of nanoparticles, although the ionic conductivity increased. One possible explanation offered was that nanoparticles cause the Li–H distance to decrease which causes an increase in the Li–H heteronuclear dipolar interaction as the main contributor to the linewidth, observed through 1H decoupling experiments22 that will be described later. While this explanation is reasonable, Mustarelli et al.20 with PEO8 + LiClO4 or LiN(CF3SO2)2 + nanoscale SiO2 found a decrease in the 7Li linewidth of nanocomposites. Bloise et al.19 with PEO + LiClO4 or LiBF4 + α-Al2O3 or γ-Al2O3 observed mixed behavior where γ-Al2O3 caused an increase in the linewidth while α-Al2O3 resulted in a decrease in the linewidth. Thus the behavior of nanoparticles clearly varies in different systems. Further linewidth measurements of γ-Al2O3 doped and undoped systems were undertaken for both 1H and 7Li species by Singh and Bhat52 combined with conductivity and thermal measurements (Figs 7.5 and 7.6). There was no significant change in Tg or in 1H NMR linewidths, indicating no effect of the filler on segmental motion. From the combination of the investigative techniques it was finally concluded that the increase in ionic motion upon
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7.5 1H linewidth vs. temperature plot for (PEG)46LiClO4 + γ-Al2O3 (10 mol%); circles and triangles represent the linewidths for the undoped and doped samples respectively. Inset shows 1H signals at room temperature. From: Th. J. Singh, S. V. Bhat, Journal of Power Sources 2004, 129, 280–7.52
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7.6 7Li linewidth vs. temperature plot for (PEG)46LiClO4 + γ-Al2O3 (10 mol%); circles and triangle represent the linewidths for the undoped and doped samples respectively. Inset (a) shows 7Li signals at room temperature. Inset (b) is an expanded plot of main figure to show sharpness of the transition for doped sample. From: Th. J. Singh, S. V. Bhat, Journal of Power Sources 2004, 129, 280–7.52
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(a) LiIn
Intensity
LiIon pair
LiOut
(b)
LiPEO LiIn
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LiOut 0 –2 –4 Chemical shift (ppm)
–6
7.7 Least square peak fits of 7Li NMR spectra of (a) SBA-15:LiClO4 (1:1 wt%); and (b) 8 wt% of SBA-15 in PEO:LiClO4. From: M. J. Reddy, P. P. Chu, U. V. Subba Rao, Journal of Power Sources 2006, 158, 614–19.55
addition of filler causes a decrease in degree of crystallinity of the polymer which is furthermore highly dependent on the amount of the filler added. Interactions in composite PEO systems with added mesoporous silica were explored by Reddy and co-workers53–55 and three possible environments of Li ions were identified by NMR studies, as shown in Fig. 7.7. In Fig. 7(a), the NMR spectrum of a Li salt solution containing silica is fitted by several peaks, a downfield peak near 0.0 ppm assigned to the Li species of an undissociated salt (ion pairs), and two upfield peaks, the middle ‘LiIn’ (0.5–0.7 ppm) and higher ‘LiOut’ (1.1–1.5 ppm), which were assigned to lithium ions coordinated on the interior channel surface and onto the outer surface of the silica particles (SBA-15), respectively. The second spectrum (Fig. 7.7b) is that of a polymer electrolyte with added silica. Here the same peaks of lithium coordinated to the silica particles are found (‘LiOut’ and ‘LiIn’) in the same positions, and while the peak representing ion pairs has disappeared a new peak designed ‘LiPEO’ emerges at (0.8–1.0 ppm) and is thought to represent lithium species coordinated to amorphous PEO. It was also shown here, that when the amount of silica SBA-15 increases the components of ‘LiOut’ and ‘LiIn’ increase dramatically while LiPEO decreases, which indicates that lithium favors the association with silica over PEO.
7.4.4 Dynamics It is very important to study the motion of the mobile ionic species in polymer electrolytes as these are dynamic environments created by the
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polymer motion in the amorphous phase and the ionic transport is closely connected to the polymer segmental dynamics. In order to investigate the dynamics of an electrolyte system, the line width of the central 7Li transition as a function of P(EO-EC)/LiCF3SO3 electrolyte content (P(EO-EC) = poly(ethylene oxide-co-ethylene carbonate)) as well as temperature were measured by Jeon and Kwak56 on solventfree polymer electrolyte materials based on poly(vinylidene fluoride) (PVdF) membranes filled with P(EO-EC). As seen in Fig. 7.8 a typical temperature dependence is observed for this system with broad spectral lines below Tg and narrow linewidths above Tg, indicating Li ion mobilities closely associated with segmental motion of P(EO–EC) polymer. In a subsequent study57 it was demonstrated how some important parameters can be further calculated from the linewidth measurements. Based on the Bloembergen–Purcell–Pound (BPP) theory40 NMR motional narrowing of the linewidth occurs when the rate of fluctuations of the local magnetic fields or electric field gradients, generally described as the correlation time, τc, is of the order of the rigid lattice line width, δωo, as described by following equation, where the measured linewidth, δυ, obeys the relation
( Δυ ) 2 =
(π )δω 2
0
2
tan −1(τ c Δυ )
[7.1]
It is generally assumed that temperature dependence of τc follows Arrhenius behavior, and then the activation energy Ea for Li ion mobility is determined by
τ c = τ 0 e( Ea / RT )
[7.2]
where τ0 is the prefactor (dwell time) corresponding to the reciprocal frequency attempt of cation jumps. From this equation both τ0 and Ea can be obtained from plotting ln τc as a function of inverse temperature, resulting in a graph such as in Fig. 7.9. The plots show two different temperature regions with widely different activation energies, separated by Tsc (temperature of slope change) that correspond to an abrupt change in ionic mobility, together with change in correlation time, activation energy, and ionic conductivity. These results are comparable to those of other polymer electrolytes such as poly(vinyl alcohol)/LiCF3SO3-based solid polymer electrolytes (SPEs)58 and polyurethane/poly-(dimethylsiloxane)/LiClO4-based SPEs.59 Similar observations have been made on conductivity measurements, where Tsc corresponds to Tg, the glass transition temperature of the amorphous polymer phase. The activation energies obtained by conductivity measurements corresponding to an energy barrier of Li ion motion are here approximately twice larger than those from 7Li NMR data. Bishop and Bray60 have also observed differences in NMR and conductivity activation energies in
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–15 3.0
3.0 (b)
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4.0
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1000/T (K–1)
7.8 (a) Temperature dependence of the 7Li NMR spectrum of E-V6E4 for a salt concentration of 1.5 mmol LiCF3SO3/g-P(EO-EC) as a function of temperature. (b) Linewidths of the 7Li NMR spectra of the three polymer electrolytes as functions of temperature. The graph in the inset shows the temperature dependence of ln τc determined from the linewidth data for E-V6E4. From: J-D. Jeon, S-Y. Kwak, Journal of Membrane Science 2006, 286, 15–21.56
lithium borate glasses because the 7Li NMR and conductivity can detect ion motions on a local scale and a more global scale, respectively. That is, the energy barrier for the Li ion to exceed in order to undergo long-range transport is higher than that of local motions.
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–12
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–14 Tsc –15 3.0
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4.0
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–1
1000/T (K )
7.9 Temperature dependence of the corresponding correlation time, τc, calculated from the 7Li NMR line width data by using the BPP equation for polymer electrolytes. From: J-D. Jeon and S-Y. Kwak, Macromolecules, 2006, 39, 8027–34.57
7.4.5 Amorphous phase transport One of the important contributions to electrolyte research coming from the NMR measurements is the recognition of ion conduction taking place predominantly in the amorphous phase of the electrolyte and only at temperatures above the glass transition of this phase. These results were obtained by motional narrowing studies, where the linewidth of resulting spectral bands is studied as a function of temperature.61–64 The results indicate that the ionic mobility is closely connected to the mobility of the polymer chains and a segmental motion-assisted ion-transport mechanism has been proposed and is now broadly accepted in the polymer electrolyte community.
7.4.6 Oriented systems This transport mechanism has been further studied on oriented systems, polymer electrolyte materials that had been mechanically65or magnetically66 oriented in a certain direction. NMR studies have revealed important facts about lithium–polymer complexes within these anisotropic systems. It was established that the re-ordering induced by the stretching extends down to the immediate environment of the cation, as evidenced by obtained angular-dependent linewidths. At the same time the diffusion coefficients were shown to increase along the stretched direction67 indicating that other factors contribute to the ion-transport mechanism, such as the alignment of the structural units of the polymer chains. Additional evidence for this was
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provided by the observation that polymer electrolytes of certain composition can be prepared in either amorphous or crystalline forms and that the latter exhibits higher conductivity in some cases.68 These results suggest a possible re-examination of Armand’s early model of ion transport in PEO in which cations move within the helical polymer structure.26 Although this mechanism has been discounted for some years, it may in fact yield a substantial contribution to the cation conductivity, but only when the helices are aligned. For randomly oriented polymer chains and helices, ionic transport assisted by amorphous polymer segmental motion remains as the dominant mechanism. However, over two decades of research on suppressing the crystalline phase and lowering the Tg of the amorphous phase have produced only incremental improvements in Li+ ion conductivity.
7.5
Relaxation processes
As a practical matter in NMR spectroscopy, relaxation times determine the length of the delay between signal-averaged acquisitions, but at the same time they are a property of the nucleus that reflects on its dynamic environment within a sample. Any excited magnetic moment relaxes back to equilibrium. There are generally two components of this relaxation for isotropic systems in the absence of chemical exchange: longitudinal or spin–lattice (T1) and transverse or spin–spin (T2) relaxations. The often measured rotating-frame relaxation time (T1ρ) is a modified T1 relaxation which is sensitive to dynamic processes of a frequency range intermediate between T1 and T2.
7.5.1 Spin–lattice relaxation (T1) Relaxation in the applied static field direction is characterized by spin– lattice or longitudinal relaxation time T1. It is the process of energy transfer from the excited nucleus to the surroundings or the lattice, usually meaning the neighboring molecules, that returns the nuclear spin system to thermal equilibrium. Since the first comprehensive treatise of relaxation published over 60 years ago,40 measurements of T1 as a function of temperature have provided a useful molecular level probe of dynamics because the most effective relaxation mechanisms, whether mediated by nuclear dipole– dipole, quadrupole, or electron spins (in samples containing paramagnetic centers), must have significant spectral density components close to the NMR frequency. Thus T1 is very sensitive to motional processes with a timescale of ~10−10 s. This can be seen from the original formulation of the relaxation rate involving a single process via:40 1 τc 4τ c ⎤ = K ⎡⎢ + T1 ⎣ 1 + ω 02τ c2 1 + 4ω 02τ c2 ⎦⎥
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where K depends on the nuclear interaction, ω0 is the NMR frequency, and τc is the correlation time. For measuring T1 a so-called inversion-recovery technique is often used, where the sample is subjected to sequences of 180° and 90° pulses separated by variable time τ. Jeon and Kwak56 have employed T1 measurements using 1 H solid state NMR to determine miscibility of P(VdF–HFP)/P(EO–EC)based porous membranes (HFP = hexafluoropropylene), where it was assumed that a single T1 relaxation complies with the condition that M − Mτ ⎤ τ ln ⎡⎢ e =− ⎥ T1 ⎣ 2 Me ⎦
[7.4]
where Mτ and Me are the intensities of the resonance at a given delay time, τ, and equilibrium state (τ ≥ 5T1), respectively. By plotting ln[(Me − Mτ)/ (2Me)] against τ, straight lines with different slopes were obtained corresponding to the pure and blend membranes, indicating a single component T1 relaxation behavior (Fig. 7.10). From the slopes of the plots, T1 values can be obtained. The fact that a single T1 was observed indicates that the spin diffusion process is sufficiently fast to equilibrate the relaxation times for all protons among the chemically different constituents and it was concluded that the blended membranes are completely homogeneous on the time and length scales probed by T1.
In[(Me – Mτ)/(2Me)]
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M-V10E0 M-V8E2 M-V6E4
–1
–2
–3
–4 0.0
0.5
1.0 1.5 Delay time (s)
2.0
2.5
7.10 Logarithmic plots of resonance intensity vs. delay time for porous membranes. The slope yields the proton spin–lattice relaxation time in the laboratory frame, T1. From: J-D Jeon and S-Y Kwak, Macromolecules 2006, 39, 8027–34.57
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7.5.2 Spin–spin relaxation (T2) Relaxation of the nuclei perpendicular to the static field direction is characterized by spin–spin or transverse relaxation time T2. This is an entropic process, where the coherence of spins is gradually lost without any energy exchange with the lattice. The main significance of this relaxation lies in its relationship with the homogeneous linewidths of the resulting spectral bands, Δω =
1 πT2
[7.5]
Δω being the spectral width at half-height. Inhomogeneous broadening mechanisms, for example from the nuclear quadrupole interaction, give addition contributions to the linewidth.
7.5.3 Spin–lattice relaxation in the rotating frame (T1ρ) In solid state systems the perturbation from the ideal radiofrequency (rf) pulses often causes random fluctuations of dipole–dipole interactions that hinder accurate measurements of the relaxation time. Undertaking measurements in a rotating frame, achieved by a spin-locking process, an unperturbed value of the relaxation time of T1ρ is obtained. Another benefit is that T1ρ is sensitive to motions on a somewhat longer timescale than T1. In contrast to the other relaxation processes T1ρ magnetization decay is often non-exponential, owing to complex interactions or a multitude of phases within the solid samples. In the case of quadrupolar nuclei, such as 7Li (I = 3/2), the spin–lattice relaxation is often described by a distribution of two or more exponential functions.69 Because of the relatively small quadrupole moment of 7Li, however, deviations from a single-exponential function are often hard to detect,59,63,70 in which case only one exponential function is necessary when evaluating 7Li relaxation measurements. Such was also the procedure of Jeon and Kwak57 when performing analysis of Li-ion transport behavior. Figure 7.11 (the lower graphs) shows the temperature dependence of the 7Li spin–lattice relaxation rates in the laboratory frame, T1−1, of P(VdF–HFP)/P(EO–EC)-based porous membranes, mentioned above. From the graph it is clear that above 240 K, T1−1 increase significantly with increasing temperature up to a maximum value of ca. 3 s−1 at T ≥ 360 K in all of the polymer electrolytes investigated. One of the samples, E-V6E4, displays well-defined T1−1 maximum at 360 K, which is an indication of only one motional process of the Li ions. In the case of the other samples, however, T1−1 maxima were not observed in the investigated temperature range. To gain better understanding of the Li ion mobility in polymer
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7.11 Temperature dependence of the 7Li spin–lattice relaxation rates, T1−1 and T1ρ−1, for polymer electrolytes measured at the Larmor frequency 155.4 MHz. The uncertainties in the 7Li spin-lattice relaxation rates are approximately 5%. From: J-D. Jeon and S-Y. Kwak, Macromolecules 2006, 39, 8027–34.57
electrolytes, 7Li spin–lattice relaxation time in the rotating frame, T1ρ, measurements were undertaken in the temperature range of 240–340 K. Figure 7.11 (the upper graphs) displays the Arrhenius plots of the 7Li T1ρ−1 for the same polymer electrolyte samples as before. For all, the T1ρ−1 plots versus reciprocal temperature exhibit clear maxima which indicate the most efficient relaxation; the T1ρ−1 values to the left of the maximum imply fast motion. The T1ρ−1 maxima (also called Tmax) of the polymer electrolytes appear to have a trend to shift to lower temperatures with increase of the amount of the P(EO-EC)/LiCF3SO3 electrolyte, indicating that the Li ions are more mobile. This can be correlated with the increased conductivity observed for these polymer electrolytes. As above, the correlation time for the relaxation process can be calculated, here expressed by the equation:
τc 1 3 4 2 ⎡ 2.5τ c 1.5τ c ⎤ = γ h⎢ + + T1ρ 2r 6 ⎣ 1 + ω 0 2τ c 2 1 + 4ω 0 2τ c 2 1 + 4ω 0 2τ c 2 ⎥⎦
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Activation energy Ea for the relaxation process, corresponding to the barrier height for the potential hindered motion, can also be obtained. For these systems the energy values are consistent with results obtained from ionic conductivity measurements and linewidth analysis.
7.6
Diffusion measurements
Another use of NMR technique in polymer electrolyte research is the measurement of self-diffusion coefficients. Here the possibility of measurements of different nuclei in one sample is fully used, and the diffusion coefficients of cationic and anionic species and the polymer chain can be obtained, together with the diffusion constant of the sometimes present additives, such as fillers or plasticizers. Diffusion experiments, first proposed by Carr and Purcell,71 are based on the fact that nuclei in a sample placed into a non-uniform magnetic field will experience different local magnetic fields depending on their spatial location. The sample is then subjected to a sequence of rf pulses. A 90° pulse is followed at time τ by a refocusing 180° pulse. This will produce a spinecho at time 2τ. The resonant spins will completely refocus only if the precession frequency is constant during the experiment. If the nuclei move along the magnetic gradient, the refocusing will be incomplete. The reduction in echo intensity due to incomplete refocusing will give the diffusion coefficient, expressed by the equation:72 2τ ⎤ ⎡ 2 δH ⎞ M ( 2τ ) = M ( 0) exp ⎢⎢ ⎥⎥ exp ⎢ − ⎛ γ Dτ a ⎥ ⎝ ⎠ δ 3 z ⎣ T2 ⎦ ⎦ ⎣ 2
[7.7]
T2 being the spin–spin relaxation time as discussed previously, γ the gyromagnetic ratio, δH/δz the magnetic field gradient, and D the diffusion coefficient. If τ is much smaller than T2, the decay of the spin-echo as a function of τ is dominated by the diffusion process and the effect of spin–spin relaxation needs not be taken into account. However in the case of a short T2, it is difficult to separate the two processes, relaxation and diffusion from each other.
7.6.1 Pulse gradient spin-echo (PGSE) measurements In this technique, the magnetic field gradient is switched off during the observation of the echo signal, first described by Stejskal and Tanner.73 This requires special equipment that allows both variation of gradient strength and variable duration and interval of gradient pulses. The contribution of T2 relaxation is here kept constant by the use of fixed delay τ between the rf pulses. The only unknown variable in the above equation is thus the diffusion coefficient. Also, it is possible to obtain high resolution spectra, which
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allow measurements of diffusion coefficients of different molecular species separately. However, this technique has its restrictions, namely in the strength of the gradient that can be used, typically ranging between 10 and 70 mT cm–1.
7.6.2 Fringe-field measurements Unlike PGSE, very strong gradients can be used in the fringe field method74,75 in which the expression previously given for static gradients applies. Compared with the pulse field gradient technique, the spectral line width is here broadened, causing the loss of sensitivity, but this is compensated for by the fact that very small diffusion coefficients can be measured (below 10−12 m2/s), in the strong fringe fields characteristic of superconducting NMR magnets. The experimental set-up uses the fact that the gradient field outside the center of the magnet decreases linearly. Thus placing the probe-head in an exactly chosen position outside the center of the magnet, the sample is exposed to a large static magnetic gradient in the order of 10 T/m. The transverse relaxation time T2 needs to be determined independently (in a uniform field experiment), as it cannot be otherwise separated from the diffusion echo decay process.
7.6.3 Using diffusion experiments Conductivity and ionic diffusion are related by the modified Nernst– Einstein equation:
σ calc =
Nq2( DLi + Danion ) α kT
[7.8]
where σcalc is the calculated ionic conductivity, N the number of ions per unit volume, q the charge of the ions, α the degree of dissociation, and DLi and Danion are the diffusion coefficients of cation and anion, respectively. It has been generally observed76–81 that the calculated conductivity with the measured self-diffusion coefficients of fluorine containing anions DF and the cations DLi (obtained from pulsed field gradient (PFG)-NMR) is larger than experimentally observed conductivity for reasons that will be discussed later. The effect of molecular orientation on ionic mobility in liquid crystal electrolytes was investigated by Saito et al.82 The liquid crystals in this study undergo a phase transition with change in molecular ordering. Conductivity measurements did not show any differences between the isotropic and ordered phases; for example, no change in slope in the conductivity plots. Therefore diffusion measurements on 7Li, 19F and 1H were conducted in order to determine individual mobilities of the involved species.
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50-80 °C 43 kJ mol–1 42 kJ mol–1 44 kJ mol–1 3.0
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1000 K / T
7.12 Temperature dependence of the diffusion coefficient of the cation (O), anion (Δ), and solvent species (䊐) from the nuclear species of 7Li, 19 F, and 1H, respectively, of LCE-B with 5 mol % Li-TFSI. Activation energies of diffusion are incorporated in the figure. From: Y. Saito, K. Hirai, S. Murata, Y. Kishii, K. Kii, M. Yoshio, T. Kato, J. Phys. Chem. B, 2005, 109, 11563–71.82
From the temperature dependence of the diffusion coefficients DLi and DF which follow closely together, it was observed that the two ionic species remain coupled, suggesting a low degree of salt dissociation (Fig. 7.12). The plot shows a change in slope at the phase transition temperature, indicating a change of the ion migration mechanism and therefore dependence on the degree of molecular orientation. It is characteristic that the activation energy of diffusion (Ea) above and below the phase transition point is reversed between the ionic species and the liquid crystal molecules (slope of DH). That is, Ea of the cation and anion diffusion in the nematic phase is lower than those of the isotropic phase. However, Ea of the diffusion of the liquid crystal molecules increases upon going to the nematic phase. This difference could be attributed to a fundamental difference in the migration mechanism of the dissolved small species and large molecules of the medium. The disagreement often observed between the ionic conductivity values measured directly and calculated from diffusion coefficient measurements via the Nernst–Einstein equation originates from several factors. The main one is attributed to the difference in the contributing species between the measurements. Electrical conductivity detects the motion of charged ions regardless of the type. However, the diffusion coefficient measures all mobile species, whether they are dissociated ions, charged multiplets, or
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electrically neutral ion pairs as long as they coexist and contain the same probed nucleus such as 7Li or 19F.83 Another factor is the possible difference in migration distance probed by self-diffusion and DC conductivity, often occurring in inhomogeneous media. Kao et al.,84 in polymer electrolytes containing LiClO4, compared DLi behavior with DPC, diffusion coefficient of propylene carbonate, a typical solvent in plasticized polymer electrolytes. A significant difference in temperature dependence was found and it could be concluded that there is very little interaction between lithium ions and the solvent. Correlation between ionic conductivity measurements and self-diffusion coefficients of lithium was also shown by Nicotera et al.85 on a group of poly(methylmethacrylate) (PMMA)–PVdF-based gel electrolytes. The mobility of lithium ion has been shown to strongly depend on the morphology of individual samples. Corain et al.88 have used measurements of diffusion coefficients in combination with electrochemical spectroscopy to investigate the degree of crosslinking inside a gel polymer electrolyte. Saunier et al.25 tried to compare diffusion coefficients of liquid electrolyte and macroporous PVdF matrix swollen with the same electrolyte, which was found to be too low for significant values to be obtained. Conductivity measurements showed a decrease by a factor of 3.6. Diffusion measurements on similar systems were conducted by several other groups.87–90 Diffusion measurements in connection with T1 relaxation experiments can provide an estimate of the diffusive pathlength and hence the sizes of blend heterogeneities,57 especially when it is not possible to measure diffusion coefficients directly. A useful approximation of the upper limit to the domain size in the case of proton NMR can be calculated by the equation where
= (6DT)1/2
[7.9]
where is the average diffusive path length for the effective spin diffusion, and D is the effective spin diffusion coefficient determined by the average proton–proton distance which determines the strength of the dipolar interaction. The relaxation time T can be represented by T1 or T1ρ according to the type of relaxation measurements.
7.6.4 Measurements of transport numbers In polymer electrolyte research one of the desired properties is high selective mobility of the cations (i.e. Li+). A multitude of approaches for increasing this selective mobility has been published,91–93 including decreasing crystalline content, inhibition of the anion mobility, solvating of cation in a mobile medium or introduction of inorganic fillers as referenced above.
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One way of measuring selective mobility in a material is calculating transference numbers of individual species as shown by Kalita et al.,14 where diffusion coefficients of measurable electrolyte species were obtained and transference numbers were calculated using the relation: t+ =
D+ D+ + D−
[7.10]
where t+ is the lithium transference number, and D+, D− are the cation and diffusion coefficients, respectively. Obtaining transference numbers for free lithium ions is not straight forward in any of the techniques commonly used. While impedance spectroscopy measures the combined mobility of any charged species, including H+ and charged ion multiplets, diffusion measurements of lithium will show diffusion of all lithium species, including neutral ion pairs, as discussed previously. Usually, the comparison between NMR and electrochemical data will give additional information about the examined material, its conduction mechanism, and aggregates formed.94
7.7
Magic angle spinning
The broadness of the spectral lines in solid state NMR is caused by strong interactions occurring in condensed phases, mainly represented by dipolar and quadrupolar interactions and chemical shift anisotropy. However, these interactions have the same geometrical scale factor, [3 cos2θ − 1], where θ is the angle between the principal interaction direction and the external magnetic field, and can be averaged by spinning the solid sample at the so-called magic angle of 54.7°, for which 3 cos2θ − 1 = 0. With increasing spinning speed a central band emerges at the isotropic chemical shift surrounded by sidebands, whose intensities can be used to calculate the chemical shift tensor. When the spinning speed exceeds the span of the chemical shift interaction, the sidebands become very small and a liquid-like spectrum is obtained. As the spinning speeds are extremely high, on the order of several times 104 rotations per second, the sample needs to be homogeneous and therefore it is possible to measure mainly finely pulverized samples. Many polymer electrolytes which are mechanically soft at room temperature require special handling procedures such as cryo-grinding in order to perform MAS measurements. A typical example of such an approach is investigation of nanoparticle surfaces where the formation and dynamics of lithium species is studied. The resulting high resolution spectra measured at different temperatures show a multitude of environments for the measured species, corresponding to different complexes, in some cases between which ionic exchange can occur.95–97
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7.8
Double resonance experiments
For better elucidation of the measured MAS spectra, several techniques well known to the polymer science community are often added to the classical MAS measurements.
7.8.1 Proton decoupling One well-known and often used technique employed in conjunction with MAS is proton decoupling, where the influence of protons surrounding the measured species, usually carbon or lithium, is inactivated or saturated by a 90° pulse train at the proton frequency, a so-called broad band decoupling. The resulting spectral line is then free from proton dipolar interactions and is considerably narrowed. Another advantage of this technique is that the heteronucleus directly bonded to a proton will be affected much more than other sites that are not, therefore facilitating spectral assignments. Decoupling protons from the 7Li MAS spectrum has for example enabled Liang et al.98 to study interaction of lithium ions with other constituents of a gel polymer electrolyte, namely with plasticizer propylene-carbonate (PC), acrylonitrile (AN) segments and PEO side chains (PEGMEN). Figure 7.13
Site 2
Site 1
Site 3
(d)
(c) (b) (a) 6
4
2
0 ppm
–2
–4
–6
7.13 Deconvolution of 7Li proton-decoupled MAS NMR spectra at 223 K for AP2/PC 50:50 wt% doped with (a) 0.25 mmol LiClO4 g−1, (b) 0.5 mmol LiClO4 g−1, (c) 1.0 mmol LiClO4 g−1, and (d) 2.0 mmol LiClO4 g−1 polymer. From: Y-H. Liang, C-C. Wang, C-Y. Chen, Journal of Power Sources 2008, 176, 340–6.98
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shows deconvolution of the resulting spectral lines for various compositions of the electrolyte. According to older works on similar systems,99–102 site 1 is associated with coordination between PC and Li+ ions, site 2 is attributed to coordination between oxygen atom on the ether groups and Li+ ions, while site 3 is related to coordination between C≡N groups and Li+. At a doping level of 0.25 mmol LiClO4 g−1 polymer (Fig. 7.13a), site 2 showed the highest intensity resonance peak in the system, indicating that Li+ ions are preferentially coordinated to PEO, because of the higher donor number of its oxygen atoms. Also, peak intensity for site 1 increases as a function of salt concentration. Therefore, the excess Li+ ions are shown to coordinate to PC. It was further shown that these different coordination situations correlate with both thermal properties and ionic conduction of the electrolytes.
7.8.2 Cross-polarization MAS is often used in combination with another double resonance technique called cross-polarization (CP). This technique makes use of the large magnetization of the protons (or in some cases 19F) that can be imparted to the heteronucleus, most commonly 13C, thus enhancing the sensitivity of the heteronuclear signal. This tool has made 13C NMR spectroscopy a routine endeavor without the need for isotopic enrichment. An additional benefit is that the longitudinal relaxation time of the protons, which is usually shorter than that of 13C is the parameter that determines the cycle time of the experiment. CP is achieved by the so-called Hartmann–Hahn condition for example, for 1H → 13C, which is γ1B1 = γ13B13, where γ is the gyromagnetic ratio, B is the amplitude of the rf field and the subscript refers to either 1H or 13C. The signal is significantly enhanced and the spectra are then further resolved by proton decoupling during detection.103 13 C CP/MAS NMR spectra were carried out at different LiClO4 concentrations to analyze the effect of the salt uptake on the polymeric chain of a polyether–siloxane hybrid and to study the chemical process involved.104 Figure 7.14 shows the 13C CP/MAS NMR spectra of the polymer doped with various amounts of LiClO4. Three major peaks at ca. 70 ppm, which are associated with the polyether backbone moieties are shown. The individual peaks are assigned to the differently bonded carbon atoms of the polymer. Upon addition of LiClO4 salt line broadening and the displacement to lower ppm values are observed for the band corresponding to methylene carbons in the PEO block (see the figure detail). This indicates that the presence of lithium salts causes a more broad distribution of the PEO-segment environments and/or reduces the segmental motion of the polymer chains; the latter results from the electronic interaction between
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10
20
30
50 78 76 74 72 70 ppm
80 25 20 15 10 ppm
Salt-free 90 80 70 60 50 40 30 20
ppm
7.14 13C CP/MAS spectra of EDS-LiClO4 complexes doped with various salt concentrations (O/Li+). The ether carbon (ca. 70 ppm) and methyl carbon (18 ppm) regions are enlarged as shown in the inset. From: W-J. Liang, Y-P. Chen, C-P. Wu, P-L. Kuo, J. Phys. Chem., 2005, 109, 24311–18.104
the Li+ cations and the ether oxygens in PEO segments. However, this interaction is less pronounced for the peaks of methylene carbons adjacent to ether oxygens (76 and 74 ppm) and methyl groups (18 ppm) on the polypropylene (PPO) segments until the salt concentration of O/Li+ = 10. Such behavior coincides with the earlier differential scanning calorimetry (DSC) results that the abrupt rise of the rate of Tg1 increases at the salt concentration of O/Li+ = 10. A similar experiment was used to characterize interactions between polymer and salt in LiClO4–PEO–poly(acrylonitrile) (PAN) hybrid polymer electrolyte.105 A broadening of the corresponding band in the observed spectra upon PAN addition was determined to indicate that a more complicated distribution of the PEO segment environments or a reduction of the segmental motion of the PEO chains results.
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7.8.3 Rotational echo double resonance nuclear magnetic resonance Despite the successes of the different NMR experiments described above in which many general structural and dynamic aspects of polymer electrolytes have been elucidated, so far none of these have been able to provide detailed information on internal structures of the electrolyte materials. Information about the short-range order on a length-scale of 1–2 Å is usually extracted from the chemical shift values obtained from static or MAS NMR. For a detailed account of the structural motifs on intermediate length scales (~4–8 Å) another approach is to use dipolar double resonance techniques such as REDOR (rotational echo double resonance) and its variants. With REDOR, the structural features are evaluated via the determination of internuclear distances between two nuclei I and S, measuring the heteronuclear dipolar coupling between these. Recently, such NMR experiments106–108 have been reported showing detailed local structure information of lithium in crystalline polymer electrolyte materials that were successfully compared with previously published X-ray diffraction results, for example 6 : 1 PEO : LiPF6 or 3 : 1 PEO : LiTf (LiTf = Lithium trifluoromethane sulfonate) samples.46 REDOR was also used to examine the microcrystalline domains of the 20 : 1 PEO : LiTf sample and it was revealed that these domains have the same local lithium environment as the completely crystalline 3 : 1 PEO : LiTf sample.46 Also, dipolar coupling constants were determined using a method, in which computer simulations with the so-called SIMPSON program109 were used to create a theoretical REDOR dephasing curve. Simulations involving one 13C and one 7Li were compared with the first four data points to determine the primary lithium–carbon dipolar-coupling constant. From this dipolar-coupling constant the average primary Li–C distances were calculated using equation D = (µ0 n) / 4π) [(γ Iγ S) / r 3]
[7.11]
where D is the dipolar-coupling constant, µ0 is the permittivity of free space, γ I the magnetogyric ratio of the dephasing nuclei (7Li), γ S the magnetogyric ratio of the observed nuclei (13C), and r the internuclear distance. 13 C–{7Li} REDOR experiments applied to polymer electrolytes have been conducted by Reichert et al.,110 who were able to determine carbon– lithium distances in crystalline PEO6LiPF6. In another study, Wickham et al.111 have used the REDOR approach to investigate the microscopic structure of the crystalline domains in PEO20LiTf. The interatomic distances obtained from solid state NMR were reported to be in excellent agreement with the values from X-ray diffraction.
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However, in the presence of motion, the efficacy of REDOR is severely limited: the dynamic processes average the dipolar interaction, consequently leading to biased, overestimated interatomic distances. Thus, when studying amorphous polymer electrolytes or the amorphous parts of a heterogeneous system, REDOR NMR spectroscopy must be performed at temperatures below the Tg, where all the mobility is suppressed.112
7.9
Two-dimensional methods
Two-dimensional (2D) spectra show either chemical shifts along both x- and y-axes, the so-called 2D correlated spectra, or J-coupling constants, for example, can be plotted against chemical shifts (2D J-resolved spectrum). The spectra can show the same nucleus, usually protons, or they can be heteronuclear, the second axis showing 13C, in most cases. 2D methods are based on couplings between nuclear dipoles or transfer of magnetization by chemical exchange. A wide variety of associated techniques exists and is readily used. Only a small number of these will be mentioned. Judeinstein et al.113 have conducted direct measurement of through-space NMR interactions that provide definitive evidence for spatial proximity of different species. Dipole–dipole interactions can be measured in principle between any NMR active nuclei with heteronuclear correlation experiments in the liquid or solid state.114 The dipole–dipole interactions decay quickly with the internuclear distances (r−3), and are difficult to evaluate for long-range distances and even more difficult when exchange, conformation, or motion phenomena are present. However, the measurement of the nuclear Overhauser method115 based on the dipole–dipole-induced crossrelaxation, was proposed to successfully measure intermolecular interactions116 and the formation of ion pairs.117,118 In agreement with recent studies, the pulsed field gradient enhanced inverse HOESY (heteronuclear Overhauser enhancement spectroscopy) sequence is usually preferred because it is more sensitive for isotope pairs 1H–7Li and also improves the digital resolution in the 1H crowded spectrum.119 Figure 7.15 presents the 1H–7Li HOESY 2D spectra for lithium triflate salt dissolved inside a block oligomer C5H11NHCONH(CH2CH2)11NHCO NHC5H11 which can be considered as a liquid model compound of the PEO/ SiO2 nanocomposites described in previous studies.120–122 The major correlation peak corresponds to the proximity of Li+ cations with the ethylene oxide segments of the heterogeneous solvent, which is obviously the most solvating moiety of these molecules. However, the small size of these segments allows some spatial proximity of the alkyl chains. Similar studies could also be used in heterogeneous and gelled SPE to observe the preferential solvating of the different constituents.123 This HOESY methodology was also used to study the ion-pairing between Li+
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OCH2 CH2
NH 1
H
ppm
7
Li
8.5 9.0 9.5 10.0 10.5 11.0 11.5 12.0 7
6
5
4
3
2
1 ppm
7.15 Heteronuclear Overhauser spectroscopy. 2D 1H-7Li correlation between the Li+ cation and the copolymer evidencing the preferential solvatation by ethylene oxide segments. From: P. Judeinstein, D. Reichert, E. R. deAzevedo, T. J. Bonagamba, Acta Chim. Slov., 2005, 52, 349–60.113
cation and fluorinated anions, tetrafluoroborate or hexafluoroborate (BF4− or PF6−) in molten PEO.
7.10
Exchange nuclear magnetic resonance
The exchange NMR method124 is another class of experiments, described by Judeinstein et al.,113 that can be used to study molecular motions which has the advantage of being much more model independent. The principles of these experiments rely on the orientation dependence of the NMR interactions for solid materials. The chemical shift in a solid is described by a 3 × 3 tensor, thus the NMR frequency of a 13C nucleus (for example) in a specific molecular site depends on the orientation of the molecule. This is known as chemical shift anisotropy (CSA) and, for amorphous or polycrystalline samples, the NMR spectrum defined by such interaction has a characteristic shape known as NMR powder patterns (Fig. 7.16a) when all orientations are present in the sample.125 The basic idea of an exchange
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H
H
B0
O O H 13C H H H O
(b) –10 Siloxane/PEO 10 hybrid
HH 13
C O H O H
110 90
70
50
Isotropic motions PEG (c) B A Siloxane structures Restrict dynamics
30
H 13
C H O
A
50
–55 °C
70 B
90
10 ppm
Simulation
30
ω1 (ppm)
(a)
Polymer electrolytes
tm = 200 ms
110 110 90 70 50 30 10 –10 w 2 (ppm)
90 70 50 30 10 –10 w 2 (ppm)
(d) tm = 200 ms
0.4 E (tm,δtCSA)
306
0.3
tm = 200 ms –55 °C
–75 °C
0.2 Isotropic motion
0.1 0.0
Isotropic motion
Small angle rotations (s = 7°)
Small angle rotations (s = 7°)
0 0.5 1.0 1.5 2.0 2.5 0 0.5 1.0 1.5 2.0 2.5 3.0 2tCSA (ms) 2tCSA (ms)
7.16 13C NMR of SiO2-PEO hybrid materials near the glass transition temperature: (a) powder pattern solid-state spectrum; (b) 13C 2D-exchange: off-diagonal signals correspond to carbon atoms experiencing changes of their orientation during the mixing time tm; (c) 1D-PUREX : bimodal curves evidence the simultaneous presence of small and high angle motions; (d) multiphasic dynamical character of PEO. Reproduced from E. R. deAzevedo, D. Reichert, E. L. G. Vidoto, K. Dahmouche, P. Judeinstein, T. J. Bonagamba, Chem. Mater., 2003, 15, 2070–8, by permission of The American Chemical Society.
NMR experiment is to take advantage of this orientation dependence for obtaining information about possible molecular reorientations. This is achieved by correlating the chemical shift frequencies of a nuclear spin observed at two different periods of time (t1 and t2) separated by a long mixing time tm (generally in the range 50–500 ms), during which motioninducing changes in the NMR frequency can occur. The usual way of correlating molecular motion with chemical shifts is by acquiring a 2D spectrum that contains the NMR frequencies before and after tm along the dimensions (ω1) and (ω2), respectively. Then, if no change in the molecular orientation occurs during tm (no molecular motions), the 2D spectrum is fully diagonal (ω1 = ω2). Changes in the molecular orientations will produce off-diagonal spectral intensities, and the final shape of the 2D spectrum will depend directly on the geometry of the molecular motion, as shown in Fig. 7.16(b) (left). In more complicated cases, the analysis of these 2D spectra can be performed using simple numerical simulations, allowing the extraction of the motion geometry, as in Figure 7.16(b) (right).
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As an example, 13C exchange NMR spectroscopy experiments have been performed on siloxane /PEO hybrid materials to get more accurate information about the polymer segmental dynamics.126 The 2D exchange spectrum shown in Fig. 7.16(b) (left) has two clearly distinct regions. In region A (see Fig. 7.16c), attributed to the carbons located in segments near the silica clusters, the 2D spectrum is purely diagonal, indicating that these segments are immobile or execute very small amplitude motion, within milliseconds on the second timescale. In contrast, off-diagonal intensities are observed in region B, which accounts for carbons in the PEO units, showing that a significant fraction of these segments e