liquid crystal polymers: from structures to applications

LIQUID CRYSTAL POLYMERS: FROM STRUCTURES TO APPLICATIONS

POLYMER LIQUID CRYSTAL SERIES

Edited by

D. ACIERNO

University of Salerno. Italy W. BROSTOW

University of North Texas. USA A.A. COLLYER

Sheffield Hal/am University, UK

Springer-sbm Archive Dordrecht

LIQUID CRYSTAL POLYMERS:

FROM STRUCTURES TO APPLICATIONS

Edited by

A.A. COLLYER Division of Applied Physics,

School of Science, Sheffield Hallam University,

UK

ELSEVIER APPLIED SCIENCE LONDON and NEW YORK

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© 1992 ELSEVIER SCIENCE PUBLISHERS LTD and THE SOCIETY OF MATERIALS SCIENCE, JAPAN

British Library Cataloguing in Publication Data

Liquid Crystal Polymers: From Structures to Applications I. Collyer, A.A. 547.7

ISBN 1-85166-797-0

Library of Congress Cataloging-in-Publication Data

Liquid crystal polymers: from structures to applications/edited by A.A. Collyer

p. cm. - (Polymer liquid crystal series) Includes bibliographical references and index. ISBN 1-85166-797-0 1. Polymer liquid crystals. I. Collyer, A.A.

Witold, 1934- III. Series. QD923.L568 1992 530.4'29--dc20

92-8744 CIP

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Preface

The subject of liquid crystals and their use in electronic displays and in non-linear optical systems has become of tremendous importance during the last decade; and the incorporation of liquid crystal units into polymeric materials has led to a group of new materials with diverse properties. Some of these properties have been utilized in new products and some have yet to be used. Much published work has appeared that deals with specific materials or particular applications, and it was felt that a book was needed to examine and explain the underlying principles governing the diverse properties of these liquid crystal polymers, LCPs.

The current work describes the diverse nature of LCPs, their synthesis, characterization, properties and finally their applications. It describes the manner in which liquid crystallinity or mesomorphism occurs in small molecules, monomer liquid crystals and polymer liquid crystals.

Chapter 1 gives a classification of the various ways in which the mesogens may be connected to the polymer chains. Currently, the bulk of LCP material is based on main chain or longitudinal LCPs for use in engineering applications. The side chain or comb polymers are intended for use in electronics and opto-electronic systems and as surfactants. Many other variants and possibilities exist but their properties have not yet been fully studied or used. In this respect it is hoped that the current work will indicate future possibilities as well as discussing current opinion.

v

vi Preface

Chapters 2 and 3 describe methods of characterizing the mesophases. In the former a comprehensive guide to textures is given. Chapter 4 describes the dielectric properties of LCPs. Chapters 5 to 8 inclusive deal with comb and longitudinal LCPs, lyotropic and thermotropic. In these chapters are discussed the syntheses, characterization, structures and properties of these materials. Chapter 9 deals with some applications of LCPs.

One of the important topics not covered in this text is that concerning the effects of electric and magnetic fields on LCPs, and associated optical effects. To cover this vast subject would be difficult in one chapter, and the merest of introductions is given at the end of Chapter 7; readers are urged to read the recent texts texts given at the end of this preface for a thorough coverage of this subject. However, it is intended to cover this area in a future volume.

'Polymer liquid crystals or liquid crystal polymers?', is a question that baffled the editor. Several reasons for using the one or the other were put forward by several authors. Professor Brostow, who has given a classification of different architectures, prefers to go in a natural progression from monomer liquid crystals to polymer liquid crystals. On the other hand, to call the book 'Polymer Liquid Crystals' may imply that the subject is about liquid crystals when it is really about polymers. A majority of the contributing authors preferred the term 'liquid crystal polymer', and as such the editor has taken refuge in the majority view.

This brief discussion then settles the aim of the book. It is intended for workers with a knowledge of polymeric materials but scant awareness of liquid crystallinity. It is hoped that this text will assist the reader in understanding the latter so that an appreciation of the way in which LCPs behave will be obtained; and, moreover, readers will be able to appraise for themselves the possibilities unfolded by the different structures and capabilities provided by these fascinating materials, with a view to extending their use in novel products and applications.

A.A. COLLYER

REFERENCES

1. Ciferri, A. (Ed.) LiqUid Crystallinity in Polymers: Principles and Fundamental Properties. VCH Verlagsgesellschaft, Germany, 1991.

Preface VII

2. McArdle, C.B. (Ed.), Side Chain Liquid Crystal Polymers. Chapman and Hall, New York, 1989.

3. Takeda, M. (Ed.), Applied Liquid Crystal Polymers. A special issue of J. Mol. Cryst. Liq. Cryst., 1989, vi.

4. Weiss, R.A. & Ober, c.K. (Eds), Liquid-crystalline Polymers. American Chemical Society, 1990.

Contents

Preface................................................................ v

List of Contributors. . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . xv

1. An Introduction to Liquid Crystallinity ............................. . W. BROSTOW

1.1 Introduction ........... . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . 1 1.2 A brief history of MLCs and PLCs ............................ 2 1.3 Mesophases and their types ................................... 3 1.4 Heterogeneous composites, molecular composites and PLCs .. 4 1.5 The nature of liquid crystallinity .............................. 8 1.6 Phases of liquid crystals ....................................... 11 1.7 Classification of PLCs ......................................... 14 1.8 Molecular structures, properties and packing .................. 21 1.9 From structures to applications ............................... 23

References ..................................................... 25

2. Characterization of Mesophases.................................... 31 C. NofOL

2.1 Introduction .................................................. 31 2.2 Typical molecular structures .................................. 32

2.2.1 Low-molar-mass liquid crystals (LMMLCs) ............. 34 2.2.2 Polymers ................................................ 41

IX

x Contents

2.3 Mesophases of rod-like molecules ............................. 45 2.3.l Nematics (N) ............................................ 46 2.3.2 Cholesterics (Ch) ........................................ 48 2.3.3 Smectics (S) .............................................. 48 2.3.4 Compounds with highly polar end groups ............... 53

2.4 Textures and polymorphism of LC polymers .................. 55 2.4.l Nematic textures ........................................ 55 2.4.2 Cholesteric textures ...................................... 60 2.4.3 Smectic textures ......................................... 62

2.5 Miscibility tests ............................................... 70 2.6 X-ray diffraction patterns ...................................... 76

2.6.1 X-ray diffraction patterns for powder samples ........... 77 2.6.2 X-ray diffraction patterns for oriented samples ........... 82

2.7 Conformation of LCPs as revealed by small-angle scattering methods .................................................... 87

References ..................................................... 91

3. NMR Studies of Thermotropic Polymers........................... 103 F. LAUPR~TRE

3.l Introduction .................................................. 103 3.2 NMR investigation of orientational conformational phenomena

in mesomorphic polymers ................................... 104 3.2.l Orientational order parameters .......................... 104 3.2.2 Principles of the NMR measurements ................... 105 3.2.3 NMR studies of orientational and conformational order in

longitudinal thermotropic polymers ................... 107 3.2.4 NMR studies of orientational and conformational order in

side chain thermotropic polymers ..................... 115 3.2.5 NMR studies of orientational and conformational order in

disco tic thermotropic polymers ........................ 122 3.3 NMR investigation of local dynamics in mesomorphic polymers 123

3.3.l Glass-liquid transition and secondary transitions ........ 123 3.3.2 Principles of the NMR experiments...................... 125 3.3.3 NMR studies of local dynamics in longitudinal liquid

crystal polymers ....................................... 127 3.3.4 NMR studies of local dynamics in side chain thermotropic

polymers .............................................. 133 3.4 NMR investigation of slow motions in mesomorphic polymers 137 3.5 Conclusions ................................................... 138

References ..................................................... 139

4. Dielectric Relaxation in Macromolecular Liquid Crystals........... 143 J.K. MOSCICKI

4.l Introduction .................................................. 143 4.2 Principles of dielectric relaxation spectroscopy ................ 144 4.3 Dielectric spectroscopy of liquid crystals ...................... 156

Contents Xl

4.3.1 Dielectric relaxation in the uniaxial phase ............... 163 4.3.2 Ferroelectric modes in chiral smectic C* phase .......... 171

4.4 Dielectric relaxation in polymers .............................. 174 4.4.1 Polymers in dilute solutions ............................. 177 4.4.2 Flexible polymers in bulk ................................ 179 4.4.3 Rod-like polymers in concentrated solutions............. 184

4.5 Dielectric spectroscopy of liquid crystal polymers ............. 187 4.5.1 Lyotropic polymers......... ............................. 191 4.5.2 Thermotropic polymers .................................. 195

References ............................................... 231

5. Lyotropic Side Chain Polymer Liquid Crystals...................... 237 P.J. HALL and G.J.T. TIDDY

5.1 Introduction .................................................. 237 5.2 Physical properties of surfactants .............................. 239

5.2.1 Dilute micellar solutions ................................. 239 5.2.2 Micelle size and shape ................................... 243 5.2.3 Liquid crystal formation of small-molecule surfactants ... 245 5.2.4 Techniques of characterization ........................... 250

5.3 Synthesis of lyotropic side chain polymer liquid crystals ....... 253 5.4 Phase behaviour of lyotropic side chain polymer liquid crystals 256 5.5 Polymerization in oriented monolayers and vesicles ........... 267

References ..................................................... 270

6. Lyotropic main chain liquid crystal polymers........................ 273 M.G. NORTHOLT and D.l. SIKKEMA

6.1 Introduction and synthetic aspects ............................ 273 6.1.1 Aromatic polyamides .................................... 274 6.1.2 Rigid-rod heterocyclic (ladder) polymers................. 276

6.2 Order in lyotropic polymer solutions .......................... 278 6.2.1 The modified Maier-Saupe mean field theory ............ 281 6.2.2 Flow behavior of lyotropic solutions .................... 286 6.2.3 Spinning of lyotropic solutions .......................... 290

6.3 Morphology of fibers and films ................................ 296 6.3.1 Chain conformation ..................................... 296 6.3.2 Results of X-ray and electron diffraction studies ......... 298 6.3.3 Characterization by optical and electron microscopy .... 305 6.3.4 Coagulation and structure formation .................... 310

6.4 Mechanical and thermal properties .................. 315 6.4.1 Elastic behavior ......................................... 318 6.4.2 Creep and stress relaxation .............................. 324 6.4.3 Strength of fibers and films .............................. 330 6.4.4 Compressive strength of fibers ........................... 334 6.4.5 Thermal properties ...................................... 337

References ............................................... 340

xii Contents

7. Thermotropic Side Chain Liquid Crystal Polymers.................. 349 DJ. SIMMONDS

7.1 Introduction .................................................. 349 7.1.1 Historical development and literature.................... 351 7.1.2 Scope and nomenclature ................................. 351

7.2 General structural features .................................... 353 7.2.1 Polymer backbone ....................................... 354 7.2.2 Spacer linkage ........................................... 355 7.23 Mesogen ......................... . . . . . . . . . . . . . . . . . . . . . . . 356 7.2.4 Other variables: disco tic and double systems ............ 357

7.3 Structure-property correlations ............................... 359 7.3.1 The backbone ........................................... 359 7.3.2 The flexible spacer ....................................... 365 7.3.3 The mesogen ............................................ 370

7.4 Copolymers ................................................... 377 7.4.1 Copolymers of mesogens with non-mesogens ............ 377 7.4.2 Copolymers with two different mesogens ................ 379 7.4.3 Cross-linked polymers-network LC elastomers ......... 381

7.5 Synthesis of comb and parallel systems ........................ 383 7.5.1 Addition polymerization of mesogenic monomers ........ 384 7.5.2 Ionic polymerization ..................................... 387 7.5.3 Condensation polymerization ............................ 388 7.5.4 Polymer modification reactions .......................... 389 7.5.5 Hydrosilylation of alkenes ............................... 390

7.6 Applications and applicable materials ......................... 391 7.6.1 Electro-optical applications .............................. 391 7.6.2 Liquid crystal elastomers ................................ 399 7.6.3 Chromatographic applications ........................... 400 7.6.4 Miscellaneous ........................ . . . . . . . . . . . . . . . . . . . 401

References ............................................... 402

8. Thermotropic Main Chain Liquid Crystal Polymers................. 407 W.A. MACDONALD

8.1 Introduction ....... . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . 407 8.2 Design of thermotropic main chain LCPs ..................... 408

8.2.1 Frustrated chain packing ................................ 410 8.2.2 Polymers with flexible spacers ........................... 411 8.2.3 Non-linear links ......................................... 412

8.3 Synthesis of thermotropic main chain LCPs ................... 416 8.4 Characterization and morphology of thermotropic main chain

LCPs ....................................................... 419 8.4.1 Solution and melt characterization ...................... 419 8.4.2 Morphology and structure of MCLCPs ................. 420

8.5 Properties of thermotropic main chain LCPs .................. 432 8.5.1 Rheology ................................................ 432 8.5.2 Mechanical properties ................................... 433 8.5.3 Miscellaneous properties ................................. 440

Contents xiii

8.6 Applications ................................................... 441 8.7 Conclusions ................................................... 442

References ..................................................... 443

9. Applications of LCP Materials..................................... 447 J.-F. JANSSON

9.1 Introduction ........... . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . 447 9.2 Injection-moulded products ................................... 450 9.3 Extruded LCP rods and profiles ............................... 454 9.4 Oriented sheets and films ...................................... 455 9.5 Thermoforming and blow moulding ........................... 456 9.6 Matrix materials for composites ............................... 456 9.7 Fibers ......................................................... 457 9.8 Coatings ................................... . . . . . . . . . . . . . . . . . . . 461

References ..................................................... 462

Index.................................................................. 465

list of Contributors

W. BROSTOW Center for Materials Characterization and Department of Chemistry, University of North Texas, Denton, Texas 76203-5308, USA

P.J. HALL

Unilever Research, Port Sunlight Laboratory, Quarry Road East, Bebington, Wirral, UK, L63 3JW

J.-F. JANSSON Polymeric Materials, Royal Institute of Technology, Stockholm, Sweden

F. LAUPRtTRE Laboratoire de Physico-chimie Structurale et Macromoleculaire associe au C.N.R.S., 10 rue Vauquelin, 75231 Paris Cede x 05, France

W.A. MACDONALD ICI Wilton Materials Research Centre, P.O. Box 90, Wilton, Middlesbrough, Cleveland, UK, TS6 8JE

J.K. MOSCICKI Institute of Physics, Jagiellonian University, Krakow, Poland; and Cornell University, Baker Laboratory, Ithaca, New York, USA

C. NOtL

Laboratoire de Physico-chimie Structurale et Macromoleculaire associe au C.N.R.S., 10 rue Vauquelin, 75231 Paris Cede x 05, France

xv

xvi List of Contributors

M.G. NORTHOLT Akzo Research Laboratories, P.O. Box 9300, 6800 SB Arnhem, The Netherlands

D.J. SIKKEMA Akzo Research Laboratories, P.O. Box 9300, 6800 SB Arnhem, The Netherlands

D.J. SIMMONDS Division of Chemistry, School of Science, Sheffield Hallam University, Pond Street, Sheffield, UK, Sl 1 WB

G.J.T. TmDY Unilever Research, Port Sunlight Laboratory, Quarry Road East, Bebington, Wirral, UK, L63 3JW

Chapter 1

An Introduction to Liquid Crystallinity

Witold Brostow Center for Materials Characterization and Department of Chemistry,

University of North Texas, Denton, Texas 76203-5308, USA

1.1 INTRODUCTION

To begin with, we divide liquid-crystalline materials into monomer liquid crystals (MLCs) and polymer liquid crystals (PLCs). This convenient terminology is due to Samulski,l who also specified that a compound is classified as an MLC irrespective of the fact whether it can or cannot polymerize. Not long ago we had an explosion of new applications of MLCs. We might well be entering now into a similar period with PLCs.

To acquire a certain perspective, I shall provide first a brief history of MLCs and PLCs. Then I shall discuss mesophases; liquid crystals constitute only one of three kinds of mesophases. Further, I shall compare heterogeneous (that is, ordinary) composites, molecular composites and PLCs. Then we shall go to the heart of this chapter: the nature of liquid crystallinity and its manifestations. On this basis we shall be able to survey existing and potential structures of PLCs, using a classification developed earlier. 2 •3 We shall see connections between a place in this classification and properties. Thus, this chapter provides an overview of the field. Chapters 1-3 form the first part of the book, including structures, characterization and dynamics. Subsequent chapters deal with specific properties, synthesis procedures, morphologies, processing and applications.

2 Witold Brostow

1.2 A BRIEF HISTORY OF MLCs AND PLCs

Researchers synthesizing liquid crystals, whether MLCs or PLCs, are proud of new materials they develop, and justly so. However, as in many areas of endeavor, humans were not exactly the first. In 1989 Li and Yu4 found that the middle gland of silk fibroin is liquid-crystallinenematic to be more accurate. Hence silkworms were producing LC materials many thousands of years before human beings started to imitate them.

The first scientific description of liquid crystals was provided by the Austrian botanist Friedrich Reinitzer. s This was in 1888, hence in 1988 we had celebrations of 100 years of discovery of liquid crystals. We also had celebrations in 1989, because in 1889 the German scientist Otto Lehmann6 coined the name 'liquid crystals'. Lehmann confirmed the experimental results of Reinitzer, which was important, since some people did not believe Reinitzer. However, Reinitzer was not exactly grateful; he maintained that the name 'liquid crystals' is wrong and constitutes a contradiction.

In 1900 or so Vorliinder started a research group working on liquid crystals and in 1908 published a book about them. 7 That group-at Martin Luther University of Halle-Wittenberg-has existed in Halle continuously up to today, with current contributions by Horst Sackmann, Dietrich Demus, Frank Kuschel, Horst Kresse and others (Martin Luther taught at that university for a number of years, including 1517 when he put his theses on that church door in Wittenberg).

After more than two decades of working on MLCs, in 1923 Vorliinder realized that PLCs must exist also. He asked: 8 'What happens to the molecules when one makes them longer and longer? Will the liquidcrystalline state disappear? From my experience there is no limit to that state by chain elongation, unless the substances could not melt any more without decomposition and could not be seen under a microscope. Starting from the p-oxybenzoic acid, Klepl9 and later Emil Fischer and his collaborators lO already obtained long chains, but did not realize their liquid-crystalline character.' Hermann Fischer gave to Vorliinder some samples prepared by his father Emil, and V orliinder found that they were liquid crystalline. 8 Thus, German researchers already had polymer liquid crystals in their hands in the nineteenth century. This helps one to keep a perspective when dealing with the latest developments in the PLC field.

An Introduction to Liquid Crystallinity 3

1.3 MESOPHASES AND THEIR TYPES

As noted in the introduction, liquid crystals represent only one of the possible types of mesophases. The name is an abbreviation of mesomorphic phases introduced by Friedel in 1922;11 he defined them as phases with microscopic structures between solids and ordinary isotropic liquids. More than three decades later, Kast 12 tried to characterize such phases in terms of lateral, longitudinal and steric disorder. Now, following Wunderlich and Grebowicz/ 3 we distinguish three kinds of mesophases: liquid crystals, plastic crystals and condis crystals.

To see what differences there are between meso phase types, we have first to describe positional, orientational and conformational disordering. A good way to do this is to consider melting of the homolog series of n-paraffins. When methane melts, various relative positions of its quasi-spherical molecules become possible; the melting process is said to be accompanied by positional disordering. There is a corresponding contribution to the entropy of fusion. Now take the next member of the series, ethane. When it melts, we have similar positional disordering (less uniform intermolecular distances); however, we also have orienta tiona I disordering: two ethane molecules can be parallel to each other, or perpendicular, or anything in between. A longer paraffin such as n-decane also undergoes positional and orientational changes on melting, but dominating in this case is conformational disordering, that is the acquisition of freedom of executing rotations about single bonds. Wunderlich and Grebowicz 13 give a good example: camphor contains 10 carbon atoms, the same as decane; however, the camphor molecule is nearly spherical and rigid, there are no orientational or conformational effects. Consequently, the entropy of fusion of camphor is much lower than that of ndecane.

All three types of meso phases have some degree of long-range order, similar to crystals. Similarly to isotropic liquids, they also have some degree of mobility other than segment vibrations. Specifically, liquid crystals exhibit positional disordering; plastic crystals show orientational disordering; and condis crystals-defined for the first time by Wunderlich and Grebowicz-show conformational disordering. Wunderlich and his colleagues stress l3- 15 that condis crystals are sometimes mistaken for liquid crystals. They also note that each of these three phases can also form a corresponding glassy phase.

4 Witold Brostow

1.4 HETEROGENEOUS COMPOSITES, MOLECULAR COMPOSITES AND PLCs

As we shall see again and again in this and in the following chapters, rigidity is an important property of most MLCs, and also of mesogenic groups in most PLCs. In turn, rigidity is connected to anisotropy of shapes and anisotropy of properties.

Rigidity of liquid crystals reminds us of rigid fiber composites such as glass-reinforced plastics. We know from books on materials science and engineering (see for instance Chapter 10 of Ref. 16) that in fiber composites the components perform different functions: rigid fibers carry load while a matrix distributes load. Mechanical properties of the composite are much better than those of either of its constituents. However, given large differences in the nature of fibers and the matrix, sufficient adhesion between the composite constituents is often a problem; cases of fiber pullout and delamination are well known. Problems of creep, fracture initiation and failure in fiber composites have been reviewed by Piggott I 7

and also by Jansson and Sundstrom. ls

We realize that the rigid reinforcement in traditional composite materials acts at the macroscopic level. As suggested in Ref. 19, we will call such materials heterogeneous composites. Problems inherent to their use can be eliminated in at least two ways. Helminiak, Hwang and their colleagues2o•21 developed materials which they called molecular composites (MCs): rigid chain molecules dispersed at the molecular level in flexible chain polymers. The other way is to use PLCs-in which, in most cases, each chain already contains rigid and flexible sequences connected by primary chemical bonds.

To break a primary chemical bond one needs an energy several orders of magnitude larger than that to overcome forces of adhesion. For this reason, in both MCs and PLCs the difficulties of achieving sufficient adhesion of large rigid units to flexible ones are eliminated from the start. Mechanical properties of PLCs are at least as good as those of the heterogeneous composites,22 while their processing is easier. We should not conclude from this that heterogeneous composites are due to disappear completely and will be replaced by MCs and PLCs. The new materials have a share of problems of their own. Monomers for PLC synthesis are often available in small quantities only, and therefore the polymers have high prices. This is the reason for blending PLCs with ordinary engineering polymers. One hopes to preserve good properties of the PLC in the blend, but to obtain the material at a much lower cost than pure PLe.


Naturally, we need to comprehend how rigid and flexible polymer sequences coexist within a material. An artist's vision can help in this. A Dutch artist, Maurits Cornelis Escher, was a master of packing together apparently disparate objects. See Figs. 1.1 and 1.2, which illustrate this. Rigid and flexible chains in MCs or rigid and flexible sequences in a PLC can pack together at least as well as crabs of two colors with their legs or as angels and devils. Further discussion of the connections between the art of Escher and science is given in Ref. 23.

Fig. 1.1. Crabs, artwork by Maurits C. Escher. (From Ref. 23.)

The materials discussed above do not yet make the full list of classes of polymer-based materials from the point of view of rigidity. At least the following classes should be listed:24

• Flexible polymers such as polysiloxanes, poly(vinyl ether), polyphosphazenes, or for that matter polyethylene.

• Semiflexible polymers. These include regular AB type copoly mers where A is stiff and B flexible. Cellulose derivatives and

6 W itold Brostow

Fig. 1.2. Angels and devils, artwork by Maurits C. Escher. (From Ref. 23.)

poly(p-hydroxybenzoic acid) (PHB) belong here, as does also for instance poly(p-phenylene terephtalamide) (PPT).

• Rigid polymers such as polyphenyl, a-helical polypeptides, or poly(pphenylenebenzobisthiazole) (PBT).

• Heterogeneous composites, which we discussed above. • Molecular composites, also already discussed. Strictly speakmg,

many of rigid chains and thus rigid rod constituents in molecular composites are liquid crystalline. We consider them separately from PLCs to follow the established usage, and also because molecular rigidity does not by itself automatically bring about liquid crystallinity. Nor is it true that liquid crystals must be rigid, since relatively flexible molecules of 1,2- or 1,3-diols with long alkyl chains in the hydrophobic part form liquid crystals. 25 . 26 As noted by Adams and Eby,27 molecular composites are of particular interest to the aerospace industry; this because of high stiffness, low density in comparison to traditional materials, and the capability


to be used for making mission-adaptive wings, that is wings with shapes which can be changed during flight. An advantage of molecular composites-or organic materials in general-for military applications is their radar invisibility (stealth). Finally, molecular composites have very good thermal stability, a property they share with polymer liquid crystals.

• Polymer liquid crystals and their blends. Incidentally, Crevecoeur and Groeninckx28.29 use the name in situ composites for PLC blends. They argue that a PLC originally dispersed in a flexible polymer as spheres or droplets can be elongated in adequate flow fields, providing in situ reinforcement.

If we compare, say, highly flexible siloxane [SiR1 R2 0]. with PBT, we realize that PLCs are between these two extremes. We can have a PLC with predominantly rigid sequences, but its melting point will be high and processing difficult. Close to the opposite end of the spectrum, we can have a PLC with predominantly flexible sequences, good processability but low melting point. This situation is advantageous. By manipulating the fraction of rigid sequences we can vary properties as well as processability. While high thermal stability is associated with more difficult processing, blending can help in this case also.

We need to know phase diagrams and phase structures as functions of the concentration of the rigid component in a series of PLCs built from the same consituents.3o For blends we also need to know phase diagrams and hierarchical phase structures as functions of concentration of blending components added to a given PLC; this makes possible intelligent processing and getting fairly close to properties defined in advance. 31

There is a variety of methods of locating phase transitions, including differential scanning calorimetry (DSC) and dynamic mechanical testing (DMT). Thermomechanical analysis (TMA), which is in reality determination of the linear isobaric expansivity (thermal expansion) turned out to be quite sensitive for transitions in PLC systems. 30,31 Among other things, these studies led to the recognition of existence in PLC systems of a phase called quasi-liquid.30

More 'exotic'-that is, so far, less frequently used-methods are also worth noting: dielectric relaxation;32-38 on which we have a whole chapter by 10zef Moscicki; thermally stimulated depolarization;39,40 electro-optical behavior41 (time dependence of transmitted light intensity under a low frequency electric field); thermo-optical analysis42 ,43 (temperature dependence of the transmission of light through birefringent

8 Witold Brostow

regions only); internal friction44 (determined on a substrate such as glass or platinum); and determination of magnetic susceptibility with a SQUID magnetometer.45

1.5 THE NATURE OF LIQUID CRYSTALLJNITY

The previous section gave us a certain background, since we now know something about classes of polymer-based materials other than PLCs. Moreover, it was noted that rigidity is related to the anisotropy of shapes and properties. We are now ready to tackle the nature of liquid crystallinity. Common to all liquid crystals is the fact that the molecules are oriented approximately parallel to a preferred axis in space called director. The degree of alignment is defined by the so-called order parameter or anisotropy factor s:

s = O' 5(3 cos 2 e - 1) (1.1)

Here e is the angle between the molecular axis and the director, and the bar above denotes an average for the material. It follows from eq. (1.1) that in a completely isotropic system s = 0 and in a perfectly aligned system s = 1. The presence of the director produces an isotropy of properties; for instance, refractive indices and dielectric constants become tensors instead of scalars. In an isotropic liquid, if we apply a field-such as an electric or magnetic field-the effects are small. The molecules respond individually. The energy difference between the alignments of a molecule parallel or perpendicular to the field is much less than the thermal energy represented by kT (k = Boltzmann's constant; T = thermodynamic temperature). By contrast, if we apply an equal field, even a weak one, to a liquid, the effects are much larger. Just as ordinary solid crystals, liquid crystals exhibit collective behavior. The already existing orientation is easily reinforced; the orienting effects of the field are, in the first approximation, proportional to volume of a collectively responding unit. Hence such a strong response.

In a PLC we can expect somewhat lower anisotropy factors, but we should realize that the rigid sequences impart orientation to flexible sequences between. Again Escher has provided something that can be used to illustrate a scientific point. Look at Fig. 1.3 and assume that the black lizards represent rigid PLC sequences and white animals the flexible sequences. It is easy to see that the flexible sequences necessarily become oriented to a certain extent 'whether they like it or not'.


Fig. 1.3. A checkerboard coloring of a tiling, artwork by Maurits C. Escher. (From Ref. 23.)

Liquid crystallinity can appear for more than one reason. Materials in which liquid crystalline properties are induced by the presence of a solvent are called lyotropic. If liquid crystallinity appears in definite temperature intervals, we have thermotropic liquid crystals. Hsiao, Shaw and Samulski46 found that liquid crystalline properties can be also brought about by elevation of pressure; I have called such Les barotropic. 3 Their existence is not surprising, since pressure and temperature changes produce similar (although not identical) effects in terms of affecting free volume.

Delineations between the above three classes of liquid crystals are not necessarily very sharp, and in general liquid crystalline phases of 'mixed'

10 Witold Brostow

character are possible. For instance, Loffier and Finkelmann47 synthesized and studied some lyotropic one-comb (see PLC classification in Section 1.7) amphiphiles (polysurfactants) which they crosslinked. The resulting swollen elastomers turned out to be liquid-crystalline also. The PLC networks had lower temperatures of transformation from liquid crystalline to isotropic than individual combs. X-ray scattering from dry and swollen elastomer samples showed that dry ones were isotropic and hence not liquid crystalline. Since water was necessary to produce liquid crystallinity, both forms can be called lyotropic. The results of Loffier and Finkelmann show that lyotropic PLCs can exist in solutions as well as swollen networks.

While we have a fair amount of knowledge about structures of liquid crystalline phases, it is still not quite clear what causes liquid crystallinity. Anisotropy of molecular shape is an important factor. Pioneering work on dilute solutions of highly asymmetrical molecules was done by Onsager48 and Isihara.49 In 1956 FlorySO pointed out that the configurational dimensions of polymer molecules in dilute solutions are often about twice those calculated assuming free rotation around all single bonds. He explained this by semi-flexibility or partial rigidity and developed an ingenious method of placing rigid highly asymmetric (rod-like) molecules on a lattice. This made possible development of a statistical mechanics of solutions of rods s1 much simpler than that of Onsager or Isihara. In 1959-60 Maier and SaupeS2 ,S3 developed a theory of nematic systems, which are the simplest liquid-crystalline ones (see the next section). There is a series of papers by Flory on liquid crystalline systems, many of them collected together in Ref. 54. We now know that, in addition to anisotropy of molecular shape, the influences of that anisotropy on molecular packing and additional stability of liquid crystalline states due to anisotropy of the dispersion forces are factors affecting liquid crystallinity.

With the importance of anisotropy well established, Krigbaum, Brelsford and Cifferiss studied the temperature variation of the axial ratio 2q/d, where q is the persistence length and d the chain diameter (the persistence length is the average sum of the projections of all bonds j ~ i on bond i in an indefinitely long chain). They found a large variation; the axial ratio falls with increasing temperature faster than was expected. Thus, a critical value is reached and a liquid crystal becomes an isotropic liquid. A different theoretical approach is that of Picken, which Northolt and Sikkema describe in Chapter 6 of this book (Section 6.2.1).


There is also experimental work aimed at understanding liquid crystallinity. Aharoni56 studied some 90 poly(ester amides) (PEA). He concluded that mesomorphism appears to be limited to highly regular, strictly alternating aromatic-aliphatic PEAs in which the methylene sequences are neither too short nor too long and where interchain hydrogen bonds hold the structure together. However, Weissflog and Demus57 synthesized 2-substituted hydroquinone bisbenzoates with large 2-substituents containing aromatic and other ring systems. Contrary to a widely held opinion, the large substituents did cause considerable deviations from the rod-like shape of the molecules, but did not prevent liquid crystallinity. Nematic and smectic phases were formed.

1.6 PHASES OF LIQUID CRYSTALS

Liquid crystals form certain characteristic phases, and this applies to both MLCs and PLCs. We shall now define the most important kinds of such phases. More information about some of these phases is provided by Claudine Noel in Section 2.3.

In Section 1.4 we defined the director. If orientation along a director is the only kind of long-range order present, we have a nematic liquid crystal. A schematic is shown in Fig. 1.4. A pile of nematic layers with the director changing from one layer to another forms a cholesteric phase. A schematic is shown in Fig. 1.5. These were the phases studied first by Reinitzer and Lehmann. The behavior of cholesterics in electric field-both MLCs and PLCs-has been reviewed by Shibaev and his colleagues. 58

• ...

• • •• • • •

Fig. 1.4. A schematic of a nematic Jiq uid crystal.

12 Witold Brostow

Fig. 1.5. A schematic of a cholesteric liquid crystal.

Further, we have a variety of smectic phases with layer structures, and some with additional long-range order within each layer. Thus, each such phase has, in addition to the director, at least one more element of a long-range order. For instance, in each smectic A phase the molecular centers lie approximately on equidistant planes perpendicular to the director. In smectic B phases we also have such planes, but additionally there is a two-dimensional hexagonal lattice within each plane. In smectic C phases there is no hexagonal structure, and the director is tilted with respect to the plane normal; the latter property distinguishes the phase C from A. There are still more. In Fig. 1.6 there are schematic representations of smectic A, B, C and G phases. The smectic G phase has herringbone symmetry.

One reason for interest in cholesteric phases is their beautiful colors. In general, liquid crystal phases have spectacular and often colorful textures, which can be seen in a book by Demus and Richter. 59 Given changes in textures and colors of cholesteric liquid crystals at specific phase transition temperatures, one of the applications of MLCs is in the field of visual temperature sensors-already used in medicine and applicable also in the automobile industry. Hence, a car part prone to overheating can be painted with a liquid crystal, and a change in color will signal that the transition temperature has been exceeded. Similarly, one paints locally the skin of a person with high fever.

Post-shearing relaxation can produce banded textures in PLCs but not in MLCs.60 A somewhat similar tractor texture was discovered by Kwiatkowski and Hinrichsen.61

Some sensors can be used to show more than one transition. This is because, as discussed by Sackmann,62 a solid upon melting sometimes produces first a more complex liquid crystalline phase, for instance a smectic one. After a further temperature increase a simpler phase, such as


.:.:.: .... :. :.',: . '.' ... '.' .. •••• :: •• ', : ••••• ,"I

Ie. ., • '., •••

(a)

(c)

(b)

//////13 \ \ \ \ \/\ / / / / /13

(d)

13

Fig. 1.6. A schematic of smectic liquid crystals: (a) A; (b) B; (c) C and (d) G phases.

smectic A and/or nematic appears. As an example, 4,4'-di-n-heptyloxyazoxybenzene goes through the following transitions:

l'd 347K . C 368K . 397K. . so I ----> smectIc -----> nematIc ----> IsotropIC

The transition at which a liquid crystal becomes isotropic is called the clearing temperature. MLC compounds, the kinds of phases they form and the respective transition temperatures are listed in books by Demus, Demus and Zaschke.63•64

An interesting phenomenon was studied by Limmer, Schmiedel, Hillner and Losche.65 In liquid crystals which produce low-temperature smectic phases (for instance G), when a high-temperature phase (smectic A) is aligned in a magnetic field, cooled down until it solidifies, and reheated again to the original temperature, the alignment is still present. Apparently, the material in the solid phase remembers what happened to it when it was still a liquid. Ukleja66 noted that a sample with a nematic and a smectic phase remembered its alignment after some years of sitting on a shelf.

14 Witold Brostow

While we said that lowering the temperature produces typically a more complex liquid crystalline phase, this is not always so. Re-entrant nematic (RN) phases, existing at temperatures below smectic phases, are also possible.67 Sackmann, Demus and collaborators68 studied binary MLC + MLC systems forming RN phases. While it was previously presumed that such phases can be formed in unary and binary systems of strongly polar compounds only, their compounds contained terminal non-polar groups. In some cases Sackmann et al. found a smectic A phase in the middle of a phase diagram, while pure components themselves form no such phases. PLCs also exhibit RN phases, as first established by Shibaev, Plate and their collaborators69 ,70 and then also by Claudine N oeI and her colleagues. 71

Phases formed and the transitions that take place are related to the problem of regularity of constituting elements of PLC chains. For instance, denoting flexible segments by F and rigid by R, we can have a regular structure ... FRFRFR ... or irregular structures. Block and graft PLCs can be made via group transfer polymerization. 72 Stupp and collaborators 73-75 generated irregular chains containing three kinds of units experimentally as well as on a computer. Stupp 7 5 introduces a concept of polyfiexibility, which implies a distribution of persistence lengths (the persistence length was defined in Section 1.5). As expected, polyflexibility affects properties, including the sharpness of the nematicto-isotropic transition; high irregularity produces a large biphasic region containing both a liquid crystal and an isotropic liquid. A Landau-type theory of Frederickson and Leibler 76 explains the experimental results of Stupp et aI., predicting that the size of the biphasic window is inversely proportional to the square root of the molecular mass.

1.7 CLASSIFICATION OF PLCs

Molecules of liquid crystals can be built from a variety of structural elements. For instance, in 1927 Vorlander 77 synthesized the first liquid crystal twins, now called tail-to-tail twins. A Siamese twin mesogen is a compound in which two independently mesogenic parts are joined in a single molecule. While earlier twins were ligated by carbon-containing groups such as methylene, Dehne et al. 78 have shown that sulfur, sulfinyl or sulfonyl groups can perform a similar role. Anatomy of MLCs from the point of view of possible structures was reviewed in detail by Demus. 79,80


Molecular structures of PLCs are usually characterized in terms of location of LC sequences. Thus one talks about main-chain PLCs in contrast to side-chain PLC polymers; the latter are also called comb liquid crystals. However, there are more possibilities than just these two. Several of them were listed by Krone, Reck and Ringsdorf.81 The first comprehensive classification scheme that enables a precise description and definition of a kind of PLC one is dealing with was proposed in Ref. 2. Since then synthetic chemists have contributed even more to the variety of PLC classes; Latin letters designating classes in the original classification2 have been replaced by Greek ones (this to avoid confusion with smectic phases, traditionally denoted by capital Latin letters). The scheme amplified accordingly3 is provided in Table 1.1. As will be seen in the following, differences in molecular structures cause large differences in properties. We now discuss each of the classes.

• Class ct, longitudinal liquid crystal polymers, earlier called main-chain polymers. A new name is necessary to distinguish them from the following classes {3 and '}' and the subclasses (S, (R and AI. There are numerous examples of class ct. We do not discuss them here in any detail since the entire Chapter 8 by MacDonald is devoted to thermotropic ones, while Chapter 6 by Northolt and Sikkema deals with lyotropic PLCs of the same class. Orientational and conformational order in longitudinal PLCs is discussed by Francoise Lauprete in Section 3.1.3. I would only like to note an account of the history of first polyesters in this class and of future perspectives by Jackson.82

• Class {3, orthogonal liquid crystal polymers. As class ct, they also contain liquid crystal groups in the main chain; however, these groups are here approximately perpendicular to the backbone. Two kinds of such polymers are obtained by Ringsdorf and coworkers,81 based on the siloxane chain and also polyesters. In this second series it appears that crystalline and liquid crystalline regions coexist until isotropization at the clearing temperature occurs .

• Class ,}" star PLCs. Four-member stars-that is crosses-were synthesized for the first time by Krone, Reck and Ringsdorf. 81 In distinction to polyesters in Class {3, which are monotropic, polymers in Class '}' are enantiotropic. Monotropic materials can only go from unstable to stable modifications, while enantiotropic undergo reversible transitions between different stable modifications. Other crosses were

16

Class

a.

EO

EP

ED

K

Witold Brostow

Table 1.1 Classification of liquid crystal polymers.

Structure

Q 0 0 Q I

Name English German

longitudinal longitudinal

orthogonal orthogonal

star (cross) Stern (Kreuz)

soft disc

rigid disc

biegsamer Diskus

steifer Diskus

multiple Multidiskus disc

one-comb Einzelkamm

multiple comb

disc comb

inverse comb

Palisadenkamm

Multikamm

f<,ammdiskus

invertierter Kamm


Table l.l-contd.

Class Structure Name

English German

61 CSC£CS parallel parallel

62 W- biparrallel biparallel

A1 --c:=--fr-=:>- mixed gemischt

A2 ~

A3 -J-{}o-

41 1 GO double doppell

412 -rn-a 5 network Netwerk

'" ( /" w }9~· conic Kegelformig

18 Witold Brostow

also synthesized; an interesting study of effects of the length of the flexible spacer in the backbone on dielectric properties is due to Kresse and his colleagues.36

• Class (, discotic polymers. Here a variety of molecules have already been synthesized, including polysiloxanes, polyamides and polyesters. These materials exhibit low molecular mobility. Two subclasses were distinguished in the original classification,2 but now we need to recognize three:3

Subclass (S, with single discs in the main chain and soft spacers between them. Two packing structures possible here and expected to coexist with each other were proposed by Wenz83 and are shown in Fig. 1.7.

(a) (b)

Fig. 1.7. Packing of ~S (disco tic soft) polymers after Wenz51 . (a) Ties leading to other stacks; (b) spacers within a stack.

Subclass (R, also with single discs in the main chain but with rigid spacers. Wendorff, Ringsdorf and collaborators84- 86 have proposed a sanidic (from the Greek for board-like) packing structure which is shown in Fig. 1.8. Subclass (M, with a multiple disc in the center. Lattermann87 has obtained two instructive examples of this subclass. One has a central core, either a benzene or a cyclohexane ring, connected via ester groups to three further cores, and on each of the latter there are three decyloxy chains in positions 3, 4 and 5. Given the length of the chains, these materials could be classified as MLCs, but there is no fundamental obstacle to making such PLCs with high molecular masses.


Fig. 1.8. Packing of ~R (discotic rigid) PLCs after Refs 52 and 53.

19

• Class e, combs or e-shaped structures, with mesogens in side chains. They were developed by Shibaev and Plate88 and independently by Finkelmann, Ringsdorf and Wendorff89 by the introduction of a flexible spacer between the backbone and the liquid crystal segment. Subsequently, Zhou, Zhu and Wen90 have shown that the flexible spacer is not necessary. The behavior of combs in mechanical, electric and magnetic fields was reviewed by Shibaev.91 We have to distinguish three subclasses, as a function of the arrangement of the side chains:

Subclass eO, combs with one row of side chains. Hardouin and his colleagues92 call these 'side-end-fixed polymers'. Usually there is a flexible tail beyond the mesogen. Sometimes there are two tails 'growing' out of the mesogen; first forked MLCs were synthesized93 .94 and later combs with forked tails.9s Many thermotropic eO PLCs are discussed in Chapter 7 by Simmonds, while lyotropic ones are covered in Chapter 5 by Hall and Tiddy. Subclass eP, combs with a palisade of side chains.96 A line drawn through the centers of gravity of the liquid crystal groups has the shape of a zigzag. Subclass eD, with multiple mesogens including double or paired ones. A number of such PLCs with polysiloxane backbones have been obtained. 97 - 100.

• Class </1, disc~comb structures. Such polymers were synthesized by Kreuder and Ringsdorf,l°l with triphenylene derivatives as the discs.

• Class K, inverse combs, that is, with rigid backbones and flexible side chains. Nematic liquid crystals were obtained by Weissflog and Demus94,102.103 with fairly short backbones. Polymers with higher molecular masses were synthesized by Ballauff.lo4.1os A number of polyesters with polystyrene side chains of different length and number per main chain were obtained by Heitz, Rohrbach and Hocker. 106

20 Witold Brostow

• Class 8, parallel structures, in which the liquid crystalline groups are in side chains and oriented approximately along the chain backbone. One can distinguish here two subclasses, depending on the location of the mesogenic groups with respect to the backbone, since apparently conformational transitions are difficult:

Subclass 81, or simple parallel. Such polymers were obtained first by Hessel and Finkelmann,l°7 later by Zhou and collaborators,l°8 and also by Keller et al.;92.109 the Bordeaux group calls them 'sideon-fixed polymers'. 92

Subclass 82, biparallel structures. This class, with nematic structures, was obtained by Hessel and Finkelmann 11 ° and called biaxial. While the word 'biaxial' is quite appropriate, it is also used in the liquid crystal field in connection with structures of smectic C phases,111.112 hence a new name was proposed. 3 A lowmolecular-mass 82 compound was synthesized by Malthete and his colleagues. 113 Properties of a polysiloxane with phenyl benzoate mesogens obtained by Finkelmann were studied by Hotz and Strobl. 114

• Class A, mixed structures, where nematogens of different shapes or orientations occur together. We have several possibilities here, including:

Subclass Ai, in which mesogens alternate between longitudinal and orthogonal positions with respect to the backbone. Such PLCs were obtained by Weissflog, Kuschel and collaboratorsYs Subclass A2, with disc-shaped and longitudinal mesogens alternating along the backbone. Subclass ,1,3, with disc-shaped and orthogonal mesogens alternating along the backbone.

• Class 1/1, double or combined PLCs, with liquid crystal groups in the main chain and in branches. The first representatives were polyesters synthesized by Reck and Ringsdorf in 1985 116 and more have been synthesized and studied since. 11 7.118 Two subclasses are possible, according to the location of side groups with respect to the backbone:

Subclass 1/11, with the spacers attached to the flexible sections of the main chain. Subclass 1/12, with the spacers attached to the meso gens in the main chain.


• Class (f, PLC networks, with elastomeric properties, obtained first by Finkelmann, Kock and Rehage119 in 1981. More such materials have since been made by Finkelmann/ 2o by Davis and Mitchell,l21 by Zentel et aU 22 ,123 and also by Jones et. al.124.125 We note that elastomers can be made from the classes listed above including longitudinal (Class IX), one-row combs (Subclass eO) and also double (Class t/I) molecules. For instance, by applying gamma radiation Shibaev, Plate and collaborators126 obtained networks from eO combs. In Table 1.1 under Class (f is shown as an example an elastomer made from comb chains. Peter and Ratzsch 127 obtained networks from longitudinal PLCs containing reactive double bonds in flexible spacers in the backbone .

• Class ill, conic molecules. Classes IX-t/I could be planar, or nearly two-dimensional. Networks are typically three-dimensional, but a planar Class (f molecule is possible, at least in principle. By contract, molecules in Class ill must be three-dimensional. Their existence was predicted by Lin 128 in 1982 and confirmed experimentally several years later.129.13o The names pyramidic or bowlic were proposed, but eventually it was decided to adopt the name conic. Lin131 predicts that these materials should have interesting electrical properties.

1.8 MOLECULAR STRUCTURES, PROPERTIES AND PACKING

After perusing the classification in the previous section, one can ask: 'It looks fine, but why do we need such a classification?' Actually, I started the development of this classification at about the same time I started to work on PLCs, simply so as not to get lost in all these structures. Analyzing the structures, I discovered another good reason:2 molecular structures of PLCs determine their phase structures and properties. This important fact is also illustrated throughout this book; hence there are separate chapters on longitudinal and on comb PLCs. Other authors have reached similar conclusions from their particular results. For instance, Ebert et al. 37 say that 'In most liquid-crystalline systems it is predominantly the molecular shape which determines which kind of liquid-crystalline phase is formed.'

In the present section I shall provide just one example of the statement just made. Consider the viscosities of two of the classes defined above. For

22 Witold Brostow

longitudinal polymers, the viscosity in the isotropic phase is typically higher than in the nematic phase. Thus, lowering the temperature so as to get below the clearing point produces a viscosity decrease. Apparently a partial alignment occurs during shear flow of the nematic. For liquid crystalline melts, however, this is not a general phenomenon. Zen tel and Wu 132

studied melt flows of simple one-row combs-that is, Subclass EO. They have found a negligible dependence of viscosities on shear rates for isotropic and nematic phases. Moreover, in contrast to longitudinal liquid crystals, viscosities in nematic phases were higher than in isotropic ones. Apparently the alignment of mesogenic side groups had-because of their flexible connections to the backbones-too weak an influence on the flow of the entire chains. In other words, even though rigid me so gens did align in nematic flow, this was not sufficient to align flexible backbones. The mechanism discussed in connection with the artwork of Escher in Fig. 1.3 was present, but for these EO PLCs it was not strong enough.

The molecular structures shown in Table 1.1 are related not only to properties and phases but also to the specifics of packing of PLC chains. Along such lines Tsukruk and his colleagues133 discuss packing in smectic phases of one-row combs as determined by X-rays. We have already shown two ways of packing of ,S discs in Fig. 1.7 and have shown sanidic packing of ,R discs in Fig. 1.8. One also talks about calamitic phases made of elongated molecules, apparently contrasting them with disc structures. Packing intermediate between the usual nematic (which is calamitic) and sanidic is also possible. 134 Destrade and his colleagues135

developed a terminology for packing structures: bacillar = finite cylinder; lenticular = disc; lamellar = infinite sheet; and columnar = infinite cylinder. Bacillar packing is expected in nematic phases and lamellar in smectic (other than smectic D) phases.

So far we have seen how the molecular structures-as represented in the classification in Table l-determine packing and phase structures. Lattermann and collaborators25 ,37,136,137 found examples of liquid crystal molecules forming aggregates. As one would expect from our earlier considerations, in such cases the shapes of the entire aggregates determine packing. Incidentally, Lattermann's work25 put to rest a widespread earlier opinion that hydrogen-bonded structures do not form liquid crystals. In fact, the aggregates of Lattermann and his colleagues are held together by hydrogen bonds. An example of the resulting columnar structure37 is shown in Fig. 1.9. In some cases a transition from hexagonal columnar to smectic structures is possible. 136 The name supramolecular mesogens has been suggested137 for liquid crystalline


Fig. 1.9. Columnar hexagonal disordered structure held together by hydrogen bonds. In the central column we see six non-disco tic tri-forked diols which

aggregate via hydrogen bonding forming a disc-like shape. (After Ref. 37.)

systems with such packing and phase structures. Hence, after assigning a new molecular structure to a PLC class, we also have to consider whether aggregate formation capabilities exist.

Computer simulations of PLC systems subjected to external mechanical forces using the method of molecular dynamics, show how mechanical properties are also affected by LC molecular shapes and packing.139 Simulation results turn out to be important for conducting intelligent processing.31

1.9 FROM STRUCTURES TO APPLICATIONS

The title of this book contains the phrase 'from structures to applications'. Any application is based on a certain property or properties. In turn, macroscopic properties are determined by an interplay of molecular structures and interactions; see Fig. 1.10. 16 There is an interesting relationship between the vertices of the triangle in Fig. 1.10: knowledge of any two vertices completely determines the third. Interactions depend on what interacts with what, and hence are related to molecular structures. We have discussed structures of PLCs in some detail, culminating in Table 1.1. To conclude this chapter, let us now make a list of essential properties of PLCs.

24 Witold Brostow

Fig. 1.10. The 'basic triangle' of materials science and engineering.

(After Ref. 16.)

Macroscopic properties

Structure Interactions

Important properties of PLCs have been summarized before2 ,3 and also discussed by Witt. 139 At least the following should be listed:

• As compared to usual engineering thermoplastics, PLCs show clear superiority with regard to chemical stability.

• PLCs show on the average lower flammability than engineering thermoplastics.

• Because of the reinforcement discussed in Section 1.4, PLCs have better moduli and overall better mechanical properties than non-PLC polymers. This shows up clearly in experiments as well as in computer modeling.

• PLCs show very low isobaric expansivity (so-called thermal expansion coefficient) and-depending on the direction-sometimes zero or negative expansivity.

• PLCs have also low isothermal compressibilities. Some longitudinal polyesters in the melt phase have compressibility of about 60% of values for non-liquid-crystal polymer melts.140

• PLCs are often easily processable with conventional processing equipment. This is an advantage not only over engineering thermoplastics, but even much more over rigid polymers and some MCs (see Section 1.4) which are soluble only in 'exotic' and highly corrosive solvents. Moreover, addition of a PLC to an engineering thermoplastic in some cases causes lowering of viscosity by several orders of magnitude. The last statement should be taken with some caution, since curves of viscosity as function of liquid crystal concent. ution can show maxima as well as minima.141-143


• For reasons discussed in the beginning of Section 1.5, PLCs orient easily with shear, electric or magnetic fields. Anisotropy can be enhanced by melt processing even at low shear rates, by application of electric fields, magnetic fields,144 drawing145 or stretching. 146 Polymer blends containing PLCs subjected to mechanical forces also acquire orientation, as shown by Springer and collaborators147 for combs blended with poly(methyl methacrylate) and polycarbonate. See also Chapter 7 by Simmonds on electric, magnetic and optical properties. We should note, however, that in some cases one needs to avoid orientation and anisotropy .

• PLCs exhibit very interesting dielectric properties; see Chapter 4 by Jozef Moscicki on this subject.

Current and potential applications of PLCs, pure or blended, depend on these properties. Some applications are not obvious from the above list. As an example, some smectic PLCs show ferroelectric and pyroelectric properties. 148 Polymer coatings for non-bake applications can be made from some eO combs.149 There are also some properties which have been observed, and which show potential for improved processing and/or improved applications. For instance, Springer and collaborators150.151 found the PLCs on the methacrylate basis aggregate in solution. Huser, Pakula and Spiess152 have studied some multiple discs, subclass ~M, cold drawing films, drawing fibers from the melt, and also achieving orientation by magnetic field imposition; mechanical properties and electrical conductivity still remain to be investigated. We conclude from all of this that PLCs should have application at least in the aerospace, aircraft, automobile, chemical, electrical and electronics industries. In this book Chaper 9 by Jan-Fredrik Jansson is devoted to applications.

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K. & Demus, D., Makromol. Chern. Rapid Commun. 1981,2,305. 42. Mucha, M. & Kryszewski, M., J. Thermal Analysis 1988, 33, 1177. 43. Mucha, M., Colloid & Polymer Sci. 1989,267,876. 44. Brostow, W. & Samatowicz, D., Polymer Eng. & Sci. 1992,33. 45. Furuya, H., Dries, T., Fuhrmann, K., Abe, A., Ballauff, M. & Fischer, E.W.,

Macromolecules 1990, 23, 4122. 46. Hsiao, B.S., Shaw, M.T. & Samulski, E.T., Macromolecules 1988,21,543. 47. Loffier, R. & Finkelmann, H., Makromol. Chern. Rapid Commun. 1990, 11,

321. 48. Onsager, L., Ann. N. Y. Acad. Sci. 1949, 51, 627. 49. Isihara, A., J. Chem. Phys. 1951, 19, 1142. 50. Flory, PJ., Proc. R. Soc. A 1956,234, 60. 51. Flory, PJ., Proc. R. Soc. A 1956,234, 73. 52. Maier, W. & Saupe, A., Z. Naturforsch. A 1959, 14,882. 53. Maier, W. & Saupe, A., Z. Naturforsch. A 1960, 15, 287. 54. Flory, PJ., Selected Works. Stanford University Press, 1985, pp. 2267-2472. 55. Krigbaum, w.R. Brelsford, G. & Cifferi, A., Macromolecules 1989,22,2487. 56. Aharoni, S.M., Macromolecules 1988,21, 1941. 57. Weissflog, W. & Demus, D., Liq. Cryst. 1988,3, 275. 58. Shibaev, V.P., Talroze, R.V., Korobeinikova, I.A. & Plate, N.A., Liq. Cryst.

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North Dartmouth, 1990. 67. Cladis, P.E., Phys. Rev. Lett. 1975, 35, 48. 68. Pelzl, G., Latif, I., Diele, S. Novak, M., Demus, D. & Sackmann, H., Mol.

Cryst. Liq. Cryst. 1986, 139,333. 69. Gubina, T.I., Kostromin, S.G., Talroze, R.V., Shibaev, V.P. & Plate, N.A.,

Vysokomol. Soed. B 1986, 28, 394.

28 Witold Brostow

70. Gubina, T.I., Kise, S., Kostromin, S.G., Talroze, R.V., Shibaev, V.P. & Plate, N.A., Liq. Cryst. 1989, 4, 197.

71. LeBarny, P., Dubois, l, Friedrich, C. & Noel, c., Polymer Bull. 1986, 15, 341.

72. Hefft, M. & Springer, l, Makromol. Chem. Rapid Commun. 1990, 11, 397. 73. Moore, lS. & Stupp, S.1. Macromolecules 1988,21, 1217. 74. Martin, P.G. & Stupp, S.I., Macromolecules 1988,21, 1222. 75. Martin, P.G., Stupp, S.1. & Moore, lS., Macromolecules 1988,21, 1228. 76. Frederickson, G.H. & Leibler, L., Macromolecules 1990,23, 531. 77. Vorliinder, D. Z. Phys. Chem. A 1927, 126,449. 78. Dehne, H., Roger, A., Demus, D., Diele, S., Kresse, H., Pe1zl, G., Wedler, W.

& Weissflog, W., Liq. Cryst. 1989,6,47. 79. Demus, D., Mol. Cryst. Liq. Cryst. 1988, 165, 45. 80. Demus, D. Liq. Cryst. 198-9,5, 75. 81. Krone, V., Reck, B. & Ringsdorf, H., presented at 16th Freiburger Arbeit-

stagung FIiissigkristalle, FreiburgjBr., 1986. 82. Jackson, W.J., Jr., Mol. Cryst. Liq. Cryst. 1989, 169,23. 83. Wenz, G., Makromol. Chem. Rapid Commun. 1985,6, 577. 84. Herrmann-Schronherr, 0., Wendorff, lH., Ringsdorf, H. & Tschirner, P.,

Makromol. Chem. Rapid Commun. 1986,7, 791. 85. Ringsdorf, H., Tschirner, P., Herrmann-Schronherr, O. & Wendorff, lH.,

Makromol. Chem. 1987,188,1431. 86. Ebert, M., Herrmann-Schonherr, 0., Wendorff, lH., Ringsdorf, H. &

Tschirner, P., Liq. Cryst. 1990, 7, 63. 87. Lattermann, G., Liq. Cryst. 1987,2, 723. 88. Shibaev, v.P. & Plate, N.A., Polymer Sci. USSR 1978, 19, 1065. 89. Finkelmann, H., Ringsdorf, H. & Wendorff, lH., Makromol. Chem, 1978,

179, 273. 90. Zhou, Q.-F., Zhu, X. & Wen, Z., Macromolecules 1989,22,493. 91. Shibaev, V. Mol. Cryst. Liq. Cryst. 1988, 155, 189. 92. Hardouin, F., Mery, S., Achard, M.F., Mauzac, M., Davidson, P. & Keller,

P., Liq. Cryst. 1990, 8, 565. 93. Nguyen, H.T., Malthete, l & Destrade, c., Mol. Cryst. Liq. Cryst. Lett.

1985, 2, 133. 94. Weissflog, w., Diele, S. & Demus, D., Mater. Chem. & Phys. 1986, 15,475. 95. Achard, M.F., Nguyen, H.T., Richard, H., Mauzac, M. & Hardouin, F., Liq.

Cryst. 1990,8, 533. 96. Duran, R., Guillon, D., Gramain, P. & Skoulios, A., Makromol. Chem.

Rapid. Commun. 1987,8, 321. 97. Engel, M., Hisgen, B., Keller, R., Kreuder, W., Reck, B., Ringsdorf, H.,

Schmidt, H.-W. & Tschirmer, P., Pure Appl. Chem. 1985,57, 1009. 98. Diele, S., Oelsner, S., Kuschel, F., Hisgen, B., Ringsdorf, H. & Zen tel, R.,

Makromol. Chem. 1987, 188, 1993. 99. Diele, S., Oelsner, S., Kuschel, F., Hisgen, B. & Ringsdorf, H., Mol. Cryst.

Liq. Cryst., 1988, 155, 399. 100. Westphal, S., Diele, S., Miidicke, A., Kuschel, F., Scheim, u., Riihlmann,

K., Hisgen, B. & Ringsdorf, H., Makromol. Chem. Rapid Commun. 1988, 9, 489.


101. Kreuder, W. & Ringsdorf, H., Makromol. Chem. Rapid Commun. 1983, 3, 807.

102. Weissflog, W. & Demus, D., Crystal Res. Technol. 1983, 18, K21. 103. Weissflog, W. & Demus, D., Crystal Res. Technol. 1984, 19, 55. 104. Ballauff, M., Makromol. Chem. Rapid Commun. 1986, 7, 407. 105. Ballauff, M. & Schmidt, G.F., Makromol. Chem. Rapid Commun. 1987,

8,93. 106. Heitz, T., Rohrbach, P. & Hocker, H., Makromol. Chem. 1989, 190, 3295. 107. Hessel, F. & Finkelmann, H., Polymer Bull. 1986, 15,349. 108. Zhou, Q.-F., Li, H.-M. & Feng, X.-D., Macromolecules 1987,20,233. 109. Keller, P., Hardouin, F., Mauzac, M. & Achard, M.F., Mol. Cryst. Liq.

Cryst. 1988, 155, 171. 110. Hessel, F. & Finkelmann, H., Polymer Bull. 1986, 15,349. 111. Bos, PJ., Pirs, J., Ukleja, P., Doane, l.W., & Neubert, M.E. Mol. Cryst. Liq.

Cryst. 1977, 40, 59. 112. Collings, P.1., Photinos, DJ., Bos, P.1. Ukleja, P. & Doane, J.w., Phys. Rev.

Lett. 1979,42, 996. 113. Malthete, J., Liebert, L., Levelut, A.M. & Galerne, C. c.r. Acad. Sci. 1986,

303, 12. 114. Hotz, W. & Strobl, G., Colloid & Polymer Sci. 1989,267, 889. 115. Rotz, u., Lindau, J., Weissflog, W., Rheinhold, G., Unseld, W. & Kuschel,

F., Mol. Cryst. Liq. Cryst. 1989,170, 185. 116. Reck, B. & Ringsdorf, H., Makromol. Chem. Rapid Commun. 1985,6, 291. 117. Endres, B.W., Ebert, M., Wendorff, J.H., Reck, B. & Ringsdorf, H., Liq.

Cryst. 1990, 7, 217. 118. Diele, S., Naumann, M., Kuschel, F., Reck, B. & Ringsdorf, H., Liq. Cryst.

1990,7, 721. 119. Finkelmann, H., Kock, M.1. & Rehage, G., Makromol. Chem. Rapid

Commun. 1981,2,317. 120. Gleim, W. & Finkelmann, H., Makromol. Chem. 1987, 188, 1489. 121. Davis, F.1. & Mitchell, G.R., Polymer Commun. 1978, 28, 9. 122. Zen tel, R. & Benalia, M., Makromol. Chem. 1987, 188,665. 123. Bualek, S. & Zen tel, R., Makromol. Chem. 1988, 189, 791, 797. 124. Dimian, A.F. & lones, F.N., Am. Chem. Soc. Symp. Ser. 1988,367, 335. 125. Wang, D. & lones, F.N., Am. Chem. Soc. Symp. Ser. 1988,367, 335. 126. Gubina, T.I., Talroze, R.V., Dakin, V.I., Imakova, N.A., Shibaev, V.P.,

Sukhov, F.F. & Plate, N.A., Vysokomol. Soed. Kr. Soobshch. 1988,30,920. 127. Peter, K. & Ratzsch, M., Makromol. Chem. 1990, 191, 1021. 128. Lin Lei, Wuli 1982,11,171; Mol. Cryst. Liq. Cryst. 1983,91,77. 129. Zimmerman, H., Poupko, R., Luz, Z. & Bilars, J., Z. Naturforsch. A 1985,40,

149. 130. Malthete, J. & Collet, A., Nouv. J. Chimie 1985,9, 151. 131. Lin Lei, Mol. Cryst. Liq. Cryst. 1987, 146,41. 132. Zentel, R. & Wu, 1., Makromol. Chem. 1986,187,1727. 133. Tsukruk, V.V., Kozlovskii, M.V., Shilov, V.V. & Shibaev, V.P., Kristallog

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30 Witold Brostow

135. Destrade, C, Foucher, P., Gasparoux, H., Nguyen, HT, Levelut, A.M. & Malthete, 1., Mol. Cryst. Liq. Cryst. 1984, 106, 121.

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1989, 47, 1841 144. Zugenmaier, P., Miigge, I In Recent Advances in Liquid Crystalline

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145. Amano, M. & Nakagawa, K., Polymer 1987,28,261 146. Zentel, R. & Benalia, M., Makromol. Chern. 1987, 188, 665. 147. Uzman, M., Kiihnpast, K. & Springer, I, Makromol. Chern. 1986, 190,

3185. 148. Patel, IS., Opt. Eng. 1987, 26, 129, 373. 149. Chen, D.-S. & Jones, F.N., J. Appl. Polymer Sci. 1988, 36, 141. 150. Springer, I & Weigelt, F.W., Makromol. Chern. 1983, 184, 1489, 2635. 151. Cackovic, H., Springer, I & Weigelt, F.W., Prog. Colloid & Polymer Sci.

1984,69, 132. 152. Hiiser, 8., Pakula, 1. & Spiess, H.W., Macromolecules 1989,22, 1960.

Chapter 2

Characterization of Mesophases

Claudine Noel Laboratoire de Physico-chimie Structurale et Macromoh!culaire associe au

C.N.R.S .. 10 rue Vauquelin. 75231 Paris Cedex 05. France

2.1 INTRODUCTION

Most crystals transform directly into the liquid phase, so that the long-range translational order and the long-range orientational order of the molecules are destroyed simultaneously. However, if the constituent molecules have pronounced anisotropy of shape, the disappearance in one, two or three dimensions of the long-range translational periodicity in the crystal may precede the collapse of the long-range orientational order. Such compounds do not show a single transition from solid to liquid but rather a cascade of transitions involving new phases; the mechanical properties and the symmetry properties of these phases are intermediate between those of a liquid and those of a crystal. These intermediate states of matter were first termed 'liquid crystals' by Lehmann (1889). They were later termed 'Les etats mesomorphes' by Friedel (1922) to avoid the ambiguities and controversies inherent in a term such as 'liquid crystals' or 'crystalline liquids' and to indicate explicitly their character as quite distinct states intermediate between the perfectly ordered periodic structure of solid crystals and the perfectly disordered structure of the amorphous state. Nowadays the terms liquid crystal (LC), mesomorphic or (mesomorphous) state and meso phase are employed almost synonymously. Substances that under suitable conditions form meso phases are referred to as 'meso gens'.

31

32 Claudine Noel

Liquid crystallinity may be induced by purely thermal effects (thermotropic systems) or by the influence of solvents (lyotropic systems). We shall be concerned here only with the former type of mesomorphism, and only in macromolecular systems, and not deal with lyotropics or low-molar-mass (LMM) systems. Mesogens embrace a diversity of structure, but with the non-amphiphilic compounds under discussion the molecules do share a common feature in being markedly geometrically anisotropic. Depending on the chemical structure and the shape of the constituent molecules or groups of molecules and on external parameters (temperature, pressure, etc.) a rich variety of phases can be observed. In this chapter we shall discuss first the molecular architectures which give rise to the known types of LCs. Then we shall give a broad classification of mesophases. It should be pointed out, however, that the detailed structures are not known with certainty for a number of phases. We shall therefore confine attention to thermotropic mesomorphism exhibited by substances based on 'lath-like' or 'rod-like' molecules. Finally, we shall discuss in detail certain of the physicochemical properties of liquid crystal polymers (LCPs) and show how studies of these particular properties can be used in elucidating the structures of the mesophases. The topics covered in this section include thermal properties, textural phenomena, miscibility tests and X-ray diffraction. Recent work by small-angle neutron scattering is also described.

2.2 TYPICAL MOLECULAR STRUCTURES

As already mentioned, the anisotropy of shape is a fundamental requirement for the formation of LCs. It is necessary to recognize a further criterion, and this relates to the rigidity of the molecule. Until recently, the rule was that the molecule has to be long and rodlike for mesomorphism to occur, but it has now been established that disc-like, pyramid-like and phasm-like (from the name of six-legged stick-like insects) molecules may also form mesophases (Fig. 2.1).1-8 The synthesis of many compounds exhibiting various kinds of LC behaviour was achieved during the last fifteen years. The reader is referred to recent reviews by Demus.9 •10 Examples of both mainchain (MC) and side-chain (SC) LCPs exist in a wide variety of the classes of high polymers.

Characterization of M esophases

Rod -like mesogens,

(i) (ii)

(K-N,118'C; N-I; 136 'c) ( K - N . 9 7.5 'C; N - [ , 122.5 'C)

(a)

Disc - like mesogens ,

R R ° I 'c"" /C""O

I ° ° 0JQr0'- /R II 0 C /C..... ° II

R ° I 0 ° ° I .. C/ C

I 0'" 'R R

(i)

Phasm -like mesogens ,

R

R

(ii)

(b)

OR OR

RO~C02 -@- 02C-@-C0 2 -®- 02C~ OR RO OR

(c)

Pyramid -like mesogens

R

R R

R R

(i) (ii)

(d)

R

R

(iii)

33

R

R

R

Fig. 2.1. Typical molecular structures. (a) Rod-like mesogens: (i) pazoxyanisole; 1 (ii) p' -n-hexyloxybenzylidene-p-n-hexyl oxyaniline. 1 (b) Discshaped mesogens: (i) hexa-substituted esters of benzene;2 (ii) hexa-substituted esters or ethers of triphenylene;3,4 (iii) hexa-n-alkanoates of truxene. s (c) Phasmlike mesogens: bis-(3,4,5-trialkoxybenzoyloxy)4' -phenyl terephthalates.6 (d) Pyramid-like mesogens: (i) hexa-substituted tribenzocyclononane; 7 (ii) octa-substituted

tetrabenzocyclododecatetraene.8

34 Claudine Noel

2.2.1 Low-molar-mass Liquid Crystals (LMMLCs)

2.2.1.1 Conventional rod-like molecules The vast majority of LMM compounds that give LCs on heating are based on 'lath-like' or 'rod-like' molecules. They are named 'calami tic'. The mesophases formed by these molecules were initially, and to a considerable extent also latterly, studied largely by organic chemists particularly interested in the relationships between the chemical constitutions of the meso gens and the types and temperature ranges of the pure mesophases produced. Reviews in Refs. 11-13 by Gray may be of value to readers who are interested in this subject.

p-Azoxyanisole (Fig. 2.1(a)) is a typical LC compound with a nearly linear structure. From a steric point of view, this is a rigid rod of length '" 20 A and width '" 5 A. The two benzene rings are nearly coplanar. The great majority of rod-like mesogenic molecules have a structure of the form

R--+-@--r; X -+-@h; R' where Rand R' represent a range of terminal substituents such as alkyl, alkoxy and cyano; a and b have small integral values and X represents a linking unit in the core structure, e.g.

-CH=N-,-N=N-,-N=N-,-COO-

~ o

In some cases, one or more phenyl rings may be saturated. To date, several thousands of mesogens are known (in the tables of Kast (1960),1 about 1400 compounds are compiled, and in the book of Demus et al. (1974),14 about 5000, with an additional 8000 in the second volume of this book (1984}.15 Typical examples are given in Table 2.1. All LC compounds with conventional rod-like molecules exhibit nematic, cholesteric, smectic and/or (in a few cases) cubic mesophases (Table 2.2).

2.2.1.2 Disc-like molecules In 1977 it was established that relatively flat, disc-shaped molecules may also form stable mesophases (Fig. 2.1).2 The great majority of discogens have a rigid central part (Figs. 2.2 and 2.3) and four to eight flexible chains (Fig. 2.4) attached to the core with a suitable linkage.

R

CnH 2n + 1-

CnH 2n + 10-

CnH 2n + 10.OC-


Table 2.1 Liquid crystals of rod-like molecules

R-@-X-@-R' x

-0. [email protected]

[email protected]

-0--CH=N-

-N=N-

-N=N-L o

-CO.O-

-CH=CH--C-C-

35

R'

-R

-F -Cl

-Br

These compounds form an entirely new class of thermotropic Les which exhibit a variety of phases, the broad structural features of which are summarized in Table 2.2. F.e. Frank has suggested that this type of mesophase may be called 'canonic' (xavwv=rod); W. Helfrich was the first to use the term columnar; the word disco tic was proposed by 1. Billard to describe the disc-like molecule as well as the mesophases. The importance of the length of the aliphatic chains suggests that the stability of the columnar phases requires (like that of the smectic) some kind of amphiphilic interaction: core and chains tend to segregate, and the chains provide lubrication between the columns. As shown in Figs. 2.2 and 2.3, the rigid core can have different symmetries but it should be noted that D2h and D4h cores do give columnar phases.

36 Claudine Noel

Table 2.2 Structural classification of thermotropic liquid crystalsa

Conventional rod-like molecules

Nematic (N)

Cholesteric (Ch)

Smectic A(SA) B(SB)

c(Scl D(SD) E(SE) F(SFl

O(SG) G'(Sa-) H(SH) H'(SH') I(Sil

Disc-like molecules

aFrom Ref. 16.

Long-range orientational order but no long-range translational order

Chiral nematic

Liquid-like layers with upright molecules Two distinct types of SB have been identified: (i) 3D crystal, hexagonal lattice, upright molecules;

(ii) stack of interacting 'hexatic' layers with in-plane shortrange positional correlation and long-range 3D six-fold 'bond-orientational' order

Tilted form of SA Cubic 3D crystal, orthorhombic, upright molecules Monoclinic (a> b) with in-plane short-range positional

correlation and weak or no interiayer positional correlation

3D crystal, monoclinic (a>b) 3D crystal, monoclinic (h>a) 3D crystal, monoclinic (a>b) 3D crystal, monoclinic (b > a) Monoclinic (b> a), possibly hexatic with slightly greater

in-plane positional correlation than SF

Columnar structure (D) with a hexagonal packing of the columns and a disordered or liquid-like arrangement of the discs in each column

Hexagonal but with an ordered arrangement of the molecular cores in each column

Rectangular, disordered

Tilted columnar structure

Nematic-like arrangement of discs

Chiral ND

Characterization of M esophases 37

R

R~~ R ('N, N,N'J

R:@:R

~/"~ R I R O~N 0 R R

T R

II

R R

R R

~ i ~

R '/_~ 7_~ R

R R R R

R R ill ill

Fig. 2.2. Discogens with hexagonal, tetragonal and trigonal thermally averaged symmetries. (From Ref. 17.)

R v

R 0 R:r9W1 R o 0 R II R

o R

R

R' R

R R'

R' R

R R'

]X

R OH-HO R 'Si / SI

R' 'OH--HO' R

1l[

R. R

'©yy©( o 0 . ,

o 6

ff~ R R X

Fig. 2.3. Discogens with binary thermally averaged symmetries. The hetero atom X can be oxygen or sulphur. (From Ref. 17.)

38 Claudine Noel

a nCmH2m+l-COO-

b nCm H2m+l -0-

c nCmH2m+l-

d nCmH2m+l -@-coo-

e nCmH2m + 1- o-@-coo-

nCI2H25-0-CO-(CH2)m -

9 nC6 H13 - CH- CH2-COO-I

CH3

F F

h nCmH2m +1- 0 * COO

F F

nCmH2m+l-V-COO- [22)

C2H5-~H-(CH2)3-0-@- COO- [23J

CH 3

k nC H2 1-0-@-m m+ "-COO-

Fig. 2.4. Side chains R (see Figs 2.2 and 2.3) used to elaborate discogens. (From Ref. 17.)

2.2.1.3 Pyramidic, cone-shaped or bowlic molecules Quite recently, in 1985, some compounds were synthesized 7 •8 •25 that are similar to discotic substances; however, in the molecules the flat core is replaced by a rigid conical unit such as tribenzocyclononene (Fig. 2.1(d)). In these compounds columnar phases with relatively high order have been found. 26

2.2.1.4 Polycatenar compounds Compounds with 2,4-disubstituted ends tanding benzenes are the bridge to substances with several endstanding chain-like substituents, now called polycatenar compounds (Table 2.3). It seems that the symmetry of the mesophases depends on the number of paraffinic chains in comparison to the core size.34 Polycatenar molecules having four to six long paraffinic chains are a good example of this idea since they present a smectic and


Table 2.3 Polycatenar compounds

CgHI90-@-cH = N-(§)-@-CN

OCgH19

Polar forked compound

C6HI3~CH = N-@--@-CN

C6H l30

13I'C . 141'C C ----+ smectlc Ad ----+ I

Biforked compounds

39

(I) [27]

(II)

[27]

Smectic C [28]

OCll~

CllH23~CH = ~oa::.-o-Coo-@-N = CH-@OCllH23 (IV)

CllH230

[29]

(V)

80'C Phasmidic phase with 82°C

C --> an oblique 2D lattice --> I [6,30]

40

Hemiphasmidic compound

C12~o--@ -c~o

C12~o-@-C~O

Swallow-tailed compound

C~SA~I

Bis-swallow-tailed compound

From Refs. 9 and to.

Claudine Noel

Table 2.3-Contd.

CO:;-@-OC12H25

86'C 69°C

I~Nb~Nl:50C 86'C ~

Cz

1

(VI)

[31]

(VII)

[32]

(VIII)

[33]

Characterization of Mesophases 41

a columnar polymorphism with a large predominance of a hexagonal mesophase for five and six chains. In this case, the column is no longer a stack of individual molecules, but of small clusters of three to five molecules. 28 Each column is still made of a central aromatic core surrounded by a paraffinic external shell but its structure is more similar to the structure of a cylinder in a hexagonal lyotropic mesophase than to that of a column made of rigid disc-like molecules. The polycatenar molecules grafted with four paraffinic chains (also called biforked mesogens) exhibit both smectic and columnar phases. In a same series a transition from a smectic C to a hexagonal phase is obtained by increasing the chain length or the temperature. Between these two phases, meso phases with two-dimensional oblique or three-dimensional cubic lattices have already been shown,29 underlining the similarities with lyotropic mesomorphism. 35

2.2.2 Polymers Traditionally, two major classes of thermotropic liquid crystalline polymers have been identified: the so called main chain (longitudinal) and side chain (comb) types (MCLCPs and SCLCPs, respectively) (Fig. 2.5). More recently other variants have appeared; these are combined LCPS36.37

which are hybrid between MCLCPs and SCLCPs, and the rigid rod types described by Watanabe et al. 38 A great wealth ofliterature already exists in the form of unified texts and reviews which detail both the major classes of LCPS.39-48 Bibliographic data have been compiled49 and reviews more or less specific to main chain 50-53 or comb54-57 polymer systems have appeared. We shall be concerned here only with main chain and side chain LCPs.

2.2.2.1 Main chain liquid crystalline polymers In the search for MCLCPs one may be inclined to select the monomers that give rise to the most rigid, extended structures. However, the requirement of a completely rigid structure must be mitigated by the requirement of polymers with accessible melting temperatures. The series of p-phenylene oligomers illustrates this point.

[email protected] When n = 2 to 4 the compounds are not rigid enough for LC formation.

At n=5 (axial ratio 5'0) a crystal- nematic (K-N) transition at 401°C is observed along with a nematic - isotropic liquid (N-I) transition at

42 Claudine Noel

Main chain lC polymers lateral substituent

X X X I I I

--;~----~~~----~~t ~----~~ I .

mesogenlc group linking group

flexible spacer

Side chain lC polymers,

polymer back bone

\

'" terminal group

Combined liquid crystal polymers,

++ Fig. 2.5. Typical structural modifications for LCPs.


445°C. When n=6 (axial ratio 5·9), the K-N transition occurs at 475°C and the N-J transition is above 600°C. When n = 7 (axial ratio 6·8), decomposition occurs below the melting point. Rigid structures such as poly( p-hydroxybenzoic acid) (PHBA) and poly(hydroquinone terephthalate) (PHQT) possess units rigid enough to show LC behaviour but their melting temperatures are close to the decomposition temperature.

-t 0 -@-CO} PHBA

t 0 -@-O-OC-@-cof PHQT The high melting behaviour of extended structures is related to

low-melting entropy. Considerable effort has been devoted to modifying the molecular architecture of rigid systems in order to reach more practical conditions for industrial processing and scientific investigation.

Four basic methods can be used to reduce the transition temperatures of MCLCPS42-45.50.53.58, and are discussed more fuily in Chapter 8:

(1) Introduction of flexible bonds or sequences to break up the rigid units and to lower the axial ratio.

(2) Introduction of kinks or bends based on essentially rigid bonds. (3) Lateral substitution of the aromatic rings to decrease molecular

symmetry, thereby disrupting the molecular packing in the crystalline state.

(4) Addition of an element of dissymmetry to the main chain by copolymerizing mesogenic units of different shapes.

These general methods are illustrated in Table 2.4. As already mentioned, mesogens embrace a diversity of structure, but

the great majority of the mesogenic groups in MCLCPs are aromatic in character and of general structure

-X-f-@--i;A-B~Y-where a and b have smail integral values.

A - B represents a linking unit in the core structure, e.g.,

-CO.O-, -N=N-,-CH=N-,-N=N-,

! ! o 0

-CH-N-, -CH-CH-, -C=C-

44 Claudine Noel

Table 2.4 General methods used for reducing the transition temperatures

Flexible sequences

Flexible bonds

-0-, -S-, -CH2-, -NH-NH-

Kinks or bends, 'Crankshaft' units

Substitution of the aromatic rings

--®- R=CnH 2n + 1 , -@ .CI.Br

R Copolymerization

-tcx:::-@-ro+ -+ o---@-o-+ -{- o--@-ro+

X and Y represent connecting units joining rigid core to flexible spacer or kink, e.g.,

-O-,-CO-O-,-O-CO-O-

2.2.2.2 Side chain liquid crystalline polymers In SCLCPs the mesogenic groups are linked directly to an eXIstmg polymer backbone or via flexible spacer (Fig. 2.5). Direct linkage59,60

except in a few cases61 - 63 gives only glasses with an anisotropy of structure that is lost at the glass transition. Coupled with steric interactions between the side groups, the tendency towards a statistical


distribution of chain conformations hinders the ordered arrangement of the pendant groups and LC formation is suppressed. Nevertheless, when the mesogenic core ordering is sufficiently strong to overcome the normal barriers associated with the random-coil conformation of the backbone, the polymer should exhibit LC properties. Decoupling of the side groups by using a flexible spacer allows the main chain motions to occur without disturbance of the anisotropic arrangement of the side chains. The polymer then may exhibit LC phases.

The most common backbones so far considered are polyacrylate, polymethacrylate and polysiloxane systems. Poly-a-chloroacrylates, itaconates, phosphazenes and ethylene oxides have also been reported. Typical spacer groups consist of between 3 and 12 methylene units. In order to enhance the degree of decoupling through a more flexible spacer, several research groups have synthesized SCLCPs containing oligoethylene64- 67 or oligosiloxane65spacers. The pendant groups in these comb-like polymers have molecular structures compatible with LC or mesophase formation as in LMM systems. The reader is referred to examples given by Demus. tO If the mesogenic side chains are rodlike (calami tic) in nature, the SCLCP may, depending upon its detailed structure, exhibit any of the accepted calami tic phases: nematic, cholesteric or smectic. Similarly, discotic side chains give disco tic phases and amphiphilic SCs give amphiphilic phases. These types are discussed fully in Chapters 5 and 7.

2.3 MESOPHASES OF ROD-LIKE MOLECULES

As already mentioned, mesogens embrace a diversity of structure. We shall not, however, be concerned here with meso phases of disc-, pyramidand phasm-like molecules, but shall confine attention to thermotropic mesomorphism exhibited by LMM compounds with elongated and relatively rigid lath-like molecules. Mesophases formed on heating such compounds are classified into three types: nematic, cholesteric, and smectic. There are more than ten recognized smectic modifications and these are denoted by SA, SB, Sc ... SL' A description of the structural features of these phases may be found in standard books and reviews 68-74 and it is my intention in this section to bring the information on this subject up to date and to present it in a more concise form.

46 Claudine Noel

2.3.1 Nematics (N) Nematic LCs can be formed by compounds that are optically inactive or by racemic mixtures.

2.3.1.1 Uniaxial nematics The nematic LC has a high degree of long-range orientational order but no long-range translational order (Fig. 2.6(a)). The preferred axis of orientation is referred to as the director and represented by an apolar unit vector n. The correlations in position between the centres of gravity of neighbouring molecules are similar to those existing in the normal isotropic liquids except for the anisotropy ~II #-~.L in the length scale.

(a)

~~-"'-r

I p

(b)

Fig. 2.6. Schematic representation of molecular arrangements in (a) nematics, and (b) cholesterics.

The direction of n is usually arbitrary in space. However, suitable treatment of the substrates between which the nematic is confined, the application of an external magnetic or electric field, and viscous flow will


result in a uniform alignment, and give a monodomain sample which is opticaliy uniaxial and strongly birefringent. Several other properties, such as the diamagnetic susceptibility and dielectric constant are also anisotropic. The anisotropy is a function of orientational order of the molecules, which decreases appreciably with rise in temperature and drops abruptly to zero at the nematic-isotropic transition in a weak first-order transition. The degree of paraliel alignment can be characterized by the order parameter:

where () is the angle between the long axis of a given molecule and the director D. The brackets indicate the average value. It should be noted that the order parameter is in general weli defined only for simplified models. The character of molecular order or disorder is usualiy not simple enough to be described and measured by a single order parameter.

2.3.1.2 Cybotactic nematics Cybotactic nematic phases show a short-range, but very definite, smecticlike ordering of the molecular centres in planes. They are the folio wing two types:

• The skewed cybotactic phase in which molecules are arranged in groups such that centres of gravity of the molecules in each group lie in a plane which makes an angle r:t. with the mean direction of the molecules in the group (Fig. 2.7(a)). This corresponds in fact to smectic C fluctuations.

• The normal cybotactic phase, which is similar to the preceding, but where r:t. is close to 90° (Fig. 2.7(b)). This corresponds to smectic A fluctuations.

(a) (b)

Fig. 2.7. Cybotactic nematics: (a) smectic C fluctuations, and (b) smectic A fluctuations.

48 Claudine Noel

2.3.2 Cholesterics (Ch) Cholesteric LCs can be formed by optically active compounds or optically active mixtures. Locally a cholesteric is very similar to a nematic material. Again, the molecular centres of gravity have no long-range order but the molecules tend to be parallel to some common axis which defines the director n. However, the structure is helicoidal with a 'precession' of the director about the helical axis Z (Fig. 2.6(b)). The structure is periodic along Z and the spatial period L is equal to one half of the pitch, P. Typical values of L are in the 3000 A range, i.e. much longer than the molecular dimensions.

2.3.3 5mectics (5) Although the total number of smectic phases is not known, the existence of the following phases seems assured 68 . 69 :

S~ S; S~ S~

(* refers to mesophases formed by optically active compounds or optically active mixtures). Smectic phases consistently exhibit higher degree of order than nematics. With the exception of SD, all smectic phases have stratified structures, but a variety of molecular arrangements are found within each stratification. The classification of SD as smectic is largely a consequence of history, and should probably be discontinued. Indeed, SD has an overall cubic symmetry and is more likely a plastic crystal.

It is useful to make a subdivision of smectics into three groups:

• The smectic phases with unstructured layers: SA and Se. • The 'hexatic' smectics: SB, SF and S). • The 'crystalline' smectics: SB, SE, SG, SH, SJ and SK. These phases

were identified as smectics on the basis of microscopic observations. However, recent high resolution X-ray studies have shown that these modifications possess three-dimensional positional order, though with extremely weak interlayer forces. These highly ordered phases cannot therefore be called LCs in the strict definition of the term.

2.3.3.1 Smectics with unstructured layers: SA, Sc and St The molecular arrangement in the smectic A phase involves a parallel arrangement of the lath-like molecules, with their ends in line, to form layers in which the long axes of the molecules tend to be orthogonal to


the layer planes, as seen in Fig. 2.8(a). Rotational motion of the molecules is fairly free, but there is no long-range regularity of packing of the centres of gravity of the molecules in the planes of the smectic layers. The layers are therefore liquid-like in nature. The movement of molecules from one layer to another occurs quite freely, and the layers are themselves quite free to slide and move over one another. The layer thickness (d) is often somewhat less than the length (L) of the molecules calculated from standard bond lengths and angles under the assumption of an all-trans conformation. A monodomain sample is optically uniaxial, the optical axis being the layer normal Z. As a rule, the smectic A phase is the less ordered smectic phase, and when a compound exhibits smectic polymorphism the smectic A phase precedes all other smectic modifications upon cooling either the isotropic liquid or the nematic (cholesteric) phase.

~ 11111/11111111111111 1 ___ _

11111/11111111111111

11111111111111111111

(a)

---- - ----~--

z

IIIIIIIIIIIII IIIIII II IIII IIIII IIIIIIII IIIIIIIIIIII III II I

(b)

-~-

/III/l1! /I//; W///// /$ ///$ 11111111/// /I/I/, 11/1/ 11/ /11/1 /1/1/ 111111111111111111 1111111111111111111111

--------

\\'\\~\\\~ M~\,\\\,\\ -------- - ---- "-,,-

\\\\\~\\\\\\~ ~\\\ \\\\\\\~\\\ -------

\\\\\ \\ \\\\\\\ ,\\\\\\\\ \\\\ \\\\\ \ \ 111111111111111111 111111111111111111111

---

Illffl III I IIIII ~ 1111/111 IIII IINI ---

/1//1/1// I! ///; Ij II l1/li ///1// (c)

Fig. 2.S. Schematic representation of molecular arrangements in (a) smectic A, (b) smectic C and (c) chiral smectic C phases.

50 Claudine Noel

The smectic C phase is the tilted analogue of the smectic A phase, i.e. the smectic layers have a liquid-like, unstructured arrangement of the lath-like molecules which are tilted with respect to the layer planes (Fig. 2.8(b)). As a consequence, the material is optically biaxial. When it occurs in the company of other phases, the smectic C phase always occurs lower in the temperature scale than either an SA or N phase, but higher in the scale than a more ordered smectic phase. This means that by heating it is possible to obtain the phase sequence: Sc ~ SA ~ N ~ I. If the smectic C phase is followed by an A phase, the tilt angle decreases gradually and finally becomes zero at the C-A transition point. If the C phase is not followed by an A phase, the tilt angle is often temperature independent and usually about 30-40°.

Smectics C can easily be twisted by the addition of optically active molecules (Fig. 2.8(c)). Pure chiral compounds showing a twisted smectic C phase (S~) have also been discovered. The structure is helicoidal with a precession of the tilt direction about the Z axis. Because of symmetry considerations, a permanent dipole moment in the plane of the chiral Sc layers is generated.75- 77 Thus S~ can exhibit ferroelectricity, as was first demonstrated by Meyer et al. 7 5 in the chiral smectic C and H phases of p-decyloxybenzylidene p'-amino-2-methylbutylcinnamate.

2.3.3.2 'Hexatic' smectics: SBHexo SF and Sf In the 'hexatic' smectic B phase, the molecules tend to be perpendicular to the layers. Each layer is again a two-dimensional liquid, but locally the molecules are distributed on a triangular lattice. The number of defects is such that the positional order does not propagate over distances larger than a few hundreds of angstroms but the bond order extends over macroscopic distances (Fig. 2.9). A monodomain sample is optically uniaxial.

It seems natural that, if one can identify phases with in-plane shortrange positional order and long-range bond order when molecules are orthogonal to the layer planes, the same should be true when the molecules are tilted. This is indeed the case. The smectic F phase is such an example. Within the layers the molecules are distributed on a nearly triangular lattice. This order extends over a few hundreds of angstroms as in the uniaxial hexatic case, but the bond ordering exhibits long-range order. The tilt points towards the side of the hexagons, which are consequently slightly distorted (Fig. 2.10(a)); this leads to a local rectangular lattice.


Fig. 2.9. Hexatic B structure. (From Ref. 69.)

(a) Tilt to side of

hexagonal net (s F)

(b) Tilt to apex of

hexagonal net (S I I

51

Fig. 2.10. (a) Short-range in-plane structure of a smectic F, and (b) short-range in-plane structure of a smectic I.

If, instead of pointing towards the side of the hexagons, the tilt points towards the apices of hexagonal net, one obtains a new phase: the smectic I phase (Fig. 2.l0(b)). Differences in the in-plane correlation length also arise, those within a layer in the S, phase being greater.

Both the SF and the S, phase are biaxial. When a material that is optically active exhibits an SF or an S, phase, then the phase itself is also optically active and is therefore called a chiral SF(S;) or a chiral S,(Sfj phase.

2.3.3.3 'Crystalline' smectics: SB, SE, SG, SH, S] and SK These phases are really crystals, albeit soft crystals. The confusion came from the fact that they are (3D) stacks of layers weakly attached to each other. Very weak forces are able to impose plastic deformations, which give some resemblance to true smectics under usual experimental conditions.

52 Claudine Noel

The smectic B phase has its constituent molecules arranged in a hexagonally close-packed array within the layers, and the molecular long axes are perpendicular to the layer planes (Fig. 2.11(a)). There is long-range hexagonal order within the layers. If the SB phase is uniaxial, the SE phase is birefringent and therefore biaxial. The biaxiality of the phase is not, however, associated with a tilted orientation of the molecules in the layers, for which X-ray studies have shown that the lamellar spacing is approximately equal to the extended molecular length (L). As shown by excellent single-crystal X-ray diffraction studies,78 it can be explained in terms of a chevron-like arrangement of the cross-sectional areas of the molecules. The SE phase has an orthorhombic symmetry (Fig. 2.11(b)).

In general, the tilted B phases are referred to as SG and SG' (or SJ) and the tilted E phases are referred to as SH and SH' (or SK) (Figs 2.11(c) and (d)).

(a) (b)

(c) (d)

Fig. 2.11. Schematic representation of molecular arrangements in (a) smectic B, (b) smectic E, (c) smectic G and (d) smectic G' (or smectic J).


2.3.4 Compounds with Highly Polar End Groups From simple energy considerations, it is evident that interactions between neighbouring dipoles can by no means be neglected in certain compounds having a strong longitudinal dipole moment (as will be the case if the molecular end group is CN or N02). In LMMLCs, this effect favours an anti parallel arrangement of the permanent dipoles. As the temperature is varied, the molecular packing is slightly altered and the resultant subtle changes in the structure appear to be responsible for a large number of smectic A and C phases and complex re-entrant behaviours.

Let us consider smectic A phases in which the molecules form a regular stack of infinite 20 liquid-like ordered layers with their director perpendicular to the layer; the presence of longitudinal dipoles raises the question of their orientation within each layer. Over the past few years69- 71 ,79,80 five SA modifications have been defined which involve different overlapping structures of the molecules; e.g. Al (monolayer), A2 (bilayer), Ad (interdigitated bilayer) and A and A (antiphase and crenelated phase, in-plane modulated structuring) (Fig. 2.12). In the Al phase the dipoles are randomly oriented in each layer, the period is then equal to the molecular length. In the SA 2 phase, the dipoles are all in the same

••

Fig. 2.12. Interdigitated bilayer (SAJ, monolayer (SAl)' bilayer SA2 and antiphase (SA) smectic A phases.

54 Claudine Noel

direction in each layer and are alternatively directed up and down in going from layer to layer. Therefore, the layer thickness is twice the molecular length. In the SAd phase, the molecules are arranged in an anti parallel, overlapping interdigitated structure with a layer spacing, d, of about 1-4 times the molecular length. The SA antiphase and the S~ crenelated phase are presumed to have an SA 2 type structure but with a long-wavelength density modulation imposed on the in-plane layer direction. The polar heads of the molecules jump periodically from one layer to the next. Analogous tilted smectic C phases are SCI' SCd , SC2 and Sc (Fig. 2.13). The Sc phase differs from the SA phase in that the lattice built by the polar heads is tilted with respect to that of the smectic layers. As a result, the optical axis is tilted with respect to both: Sc are biaxial phases. Different polymorphic forms have also been observed in the ordered SB81 •82 and SE 83- 85 phases of strongly polar compounds. In addition, some remarkable structures have been discovered which relate to recent theoretical discussions. 86 •87 The experimental evidence shows that a liquid crystal phase can support, simultaneously, two mass-density fluctuations of incommensurate wavelengths. 82.83,88-90

Normally in phase transitions the higher-temperature phase is less well ordered than the lower-temperature phase. This is not a law of thermodynamics, however, and it has been found that certain cyano-

Fig. 2.13. Sc" SCd' SC2 and Sc phases.


compounds with alkyl or alkoxy end groups exhibit re-entrant phenomena. The re-entrant nematic phase was first discovered by Cladis91 in binary mixtures of cyano-compounds. The sequence of transitions on cooling was:

I --t N --t SAd --t NR --t Solid

The lower-temperature nematic phase was referred to as re-entrant nematic phase (NR)' Examples of more complex re-entrant behaviour are now known for pure compounds.16.69-71.79.92

2.4 TEXTURES AND POLYMORPHISM OF LC POLYMERS

Texture is the LC analogue of morphology in solid crystals. Mesophases usually show various singularities in the distribution of the molecules which are characteristics of the structural features of the perfectly ordered phases. These singularities may gather, in simple experimental situations, in more or less regular sets forming LC textures. For a practical and relatively fast classification of LC, the microscopic observations of the textures are most useful. There are limitations, however, and a complete classification of smectic phases by textures is not always possible: similar textures may be observed with two LC states separated by a phase transition.

Two relevant books with photographic illustrations have been published dealing deeply with the textures of LMMLCs.68 .69 The primary aim of this section is therefore to present to the reader textures which are typical of nematic, cholesteric, smectic A and smectic C phases and to show how these types of meso phases can be identified and distinguished from one another. In this regard, however, it is important to note that microscopic observations are sometimes misleading owing to the difficulty with which LCPs give specific textures in the liquid crystalline state. This might be due to their multiphase nature (existence of polycrystalline and amorphous material), polydispersity and/or the high viscosities of the liquid crystalline melts. In most cases, samples must be annealed for hours or days at a suitable temperature if textures reminiscent of those of LMMLCs are to be seen.93 •94

2.4.1 Nematic textures When the isotropic liquid is cooled, the mesophase may appear at the clearing point in the form of separate small, round objects called droplets

56 Claudine Noel

(Fig. 2.14).95 This indicates that the polymer under investigation exhibits a nematic phase since droplets occur nowhere else. After further cooling the droplets join together to form larger structures from which stable texture finally forms.

In most cases the textures exhibited by polymers in the nematic state are similar to those of LMMLCs. The Schlieren, threaded and marbled textures are the ones most often cited in the literature. Examples of typical polymeric textures are shown in Figs 2.15 and 2.16.96 ,97

Fig. 2.14. Nematic droplets. Crossed polarizers. Original magnification x 200. (Ref. 95.)

A suitable treatment of the glass surfaces between which the polymer is observed allows one to obtain planar layers in which the director lies parallel to the surface. A number of techniques have been used to create reproducible planar alignment. Rubbing of a raw glass plate with a diamond paste leads to uniformly aligned samples.98 ,99 SiOx layers evaporated obliquely (at 60° incidence)97,loo and freshly cleaved mica surfaces101 produce similar effects. Polymer coatings on glass substrates100,102-106 and surface-deposited hexadecyltrimethylammonium bromide97 can also be used to align LCPs homogeneously. However, since a non-degenerate planar alignment is generally required, rubbing of the substrate after deposition of the polymer is used. Viewing planar samples from the top between crossed polarizers results in the observation of four positions of extinction.


Fig. 2.15. Schlieren texture. Crossed polarizers. Original magnification x 200. (Ref. 96.)

Fig. 2.16. Threaded texture. Crossed polarizers. Original magnification x 200. (Ref. 97.)

58 Claudine Noel

Preparations without special surface treatments often yield planar layers which are more or less homogeneous. These layers exhibit the Schlieren texture. When viewed between crossed polarizers this texture shows an irregular network of black brushes branching out from a number of scattered points or 'nuclei' and passing continuously from one nucleus to another. They correspond to extinction positions of the nematic liquid. Four brushes meet at some nuclei (integral nuclei) and two at others (half nuclei). The points are the intersections of vertical lines of singularity with the glass surfaces. 68.107 Possible topologies (molecular arrangements) about point singularities are given in Fig. 2.17. The strength of the disclination, S, is connected with the number of dark brushes meeting at one point:

ISI=number ofbrushes/4

From the observation of singularities with S = ± 1/2, the meso phase can be identified unambiguously as a nematic phase since these singularities occur nowhere else. Smectic C phases, which can exhibit the Schlieren texture, only form singularities with four brushes corresponding to S= ±1.

Fig. 2.17. Schematic diagram of molecular trajectories associated with dis

c1inations of strength ± ~ or ± 1.

~~ ~~

5 =-1

((~

In some preparations, the disclination lines do not lie perpendicularly, but more or less horizontally, their ends being attached to the glass surfaces and the other parts floating freely in the LC. The term 'threads' is usually used to describe these disclination lines and the corresponding texture is the threaded texture.

Some optical observations of defects in polymeric nematic phases exhibiting threaded and Schlieren textures showed that a massive organization of macromolecules occurs upon extended annealing or as the temperature is raised. 97. 108-112 Immediately after melting of the specimen, a large number of disclination lines are observed. However, the


number of disclination lines decreases with annealing time or increasing temperature. The free energy of the system is relaxed as a consequence of the annihilation of the disclinations and hence of a coarsening of the texture. Recent experiments by Percec et al. 112 showed that the rate of decrease of the number of disclinations per unit volume is strongly influenced by the molecular weight of the polymer. It decreases drastically with increasing polymer molecular weight. This behaviour must be carefully considered when studying the physical properties of polymers, since the number of disclinations per unit volume dictates the size of the nematic 'domain'.

Nematic marbled texture consists of several areas with different molecular orientation. On observing the preparation between crossed polarizers, one can note that the interference colour is nearly constant within the individual areas, indicating quasi-homogeneous regions.

Homeotropic texture caused by a spontaneous orientation of the sample is also found (Fig. 2.18).95 The field of view using crossed polarizers and orthoscopic illumination remains uniformly dark as the preparation is turned because the director is oriented perpendicular to the glass surfaces and parallel to the light beam. If the cover glass is

Fig. 2.18. The Schlieren and homeotropic textures promoted by the use of very clean surfaces. Crossed polarizers. Original magnification x 200. (Ref. 95.)

60 Claudine Noel

touched, the originally dark field of view brightens instantly, thus distinguishing between homeotropic and isotropic texture. The formation of a homeotropic texture can be promoted by the use of very clean glass surfaces95.97 or by suitable surface treatment with a lecithin solution.1 03.104,113.114 Homeotropic alignment can also be produced in experiments using surface-active molecules covalently bound to the substrate. 115

A final remark will concern the 'glassy liquid crystal'. All the textures observed for both MC and SCLCPs in the nematic state can easily be quenched and supercooled to room temperature. In spite of the thermal shock, both the planar alignment and the homeotropic one can be retained in the glassy state. 1 0 1.116-120

2.4.2 Cholesteric Textures Chiral compounds on addition to polymers in the nematic state yield cholesteric liquids.97.101,121 123 The same effect can be obtained by copolymerization (i) of a binary mixture of chiral monomers124 or (ii) of a nematogenic monomer with a chiral comonomer.122,125133 These cholesteric LCs can occur in the planar texture, the sample being uniformly aligned with the twist axis perpendicular to the plane of the film. There are, however, often alignment discontinuities which form pattern-like cracks (Fig. 2.19). 'Moire' fringes can also be observed (Fig. 2.20). In the planar texture, these cholesterics can show reflection colours. The wavelength of the light at the centre of the reflection band is, for perpendicular incidence, equal to the length of the pitch multiplied by the

7

Fig. 2.19. Cholesteric texture obtained by dissolving a chiral compound in co polyester prepared from terephthalic acid, methylhydroquinone, and pyro

catechol. (From Ref. 97.)


Fig. 2.20. Typical planar texture with moire fringes obtained by the addition of an optically active compound to the polyazomethine:

==*N~N = Crr-@-CH=F

cli:3

mean refractive index. The pitch of the helix is highly dependent on the temperature of the preparation, the nature of the organic materials comprising the mesophase, and the content of the chiral component. 54 The optical activity of a planar cholesteric layer is directly observable by the asymmetrical colour changes that appear upon rotating the analyser. In polarization microscopy, for a dextro-mesophase, a small clockwise rotation of the analyser shifts the colour to longer wavelengths.

Wedge-shaped samples, when the surfaces are prepared for a parallel alignment, show a particular feature called Grandjean steps or threads. The threads follow lines of equal thickness and are connected with a discontinuous change of the pitch. In the wedge-shaped sample, the pitch decreases with increasing thickness because the orientation at the surfaces is fixed. When the deformation energy becomes too high, a disclination line forms that allows a discontinuous change of pitch.

More recently it has been shown that it is possible to prepare cholesteric homopolymersyz.134-138 This possibility can lead to strongly twisted cholesterics which may occur in textures reminiscent of those of smectics, especially smectics A.68 Such cholesteric substances may yield non-planar fan-shaped, focal-conic or polygonal textures. In the case of homopolymers

62 Claudine Noel

prepared from chiral monomers, this makes it difficult to interpret the observed textures in terms of cholesterics or smectics. However, by adding a nematic solute, the pitch can be increased and planar textures can be observed, thus distinguishing between cholesteric and smectic phases.

2.4.3 Smectic Textures In the case of smectic polymers, observation of specific textures may be difficult. Often textures occur whose characteristics are somewhat obscure and observable only with difficulty even after extended annealing within the smectic phase. However, several variants occurl16,121.139-148 which closely resemble the focal-conic and fan-shaped textures of A and C modifications in LMMLCs.68 ,69

A focal-conic domain (Fig. 2.21)107 consists of two focal-conics, an ellipse and a branch of hyperbola. The molecules lie on the lines drawn between any point A on the ellipse (E) and any point B on the hyperbola (H). All the lines starting from A and describing (H) form a cone of revolution and similarly for the lines starting from any point on (H) and describing (E). The smectic layers are perpendicular to these lines and form a family of Dupin cyclides. The focal-conic domains are limited by the cones tapering to the limiting points on (E) and (H) and are tangential to the adjacent focal-conic domains along these cones. The focal-conic domains lead to beautiful arrangements described as polygonal and focal-conic textures (Fig. 2.22).

Smectic A phases often display textures which are probably based on the focal-conic arrangement but in which the ellipses and hyperbolae cannot always be distinguished in the microscopic image. The chief examples are, firstly, the so-called batonnets. 116,139.140,151 When the

(E)

Fig. 2.21. Scheme of focal-conic domains. Straight lines such as AB joining points on the ellipse to points on the hyperbola mark the direction of n in the case of smectics, and the direction of the helical axis in the case of cholesterics. (From

Ref. 107.)


Fig. 2.22. Simple focal-conic texture for a SA homopolyester prepared from din-propyl-p-terphenyl-4,4" carboxylate and HO- CH(CH 3)-CH(CH 3)-OH.

(From Ref. 149.)

smectic A phase is obtained directly from the isotropic liquid, elongated sharp-pointed smectic particles, the batonnets, are formed at the clearing point (Fig. 2.23). On cooling they grow, merge and may result in a focal-conic texture. Secondly, there is the fan-like texture. In thin layers, the focal-conic texture changes its appearance and gives the fan texture or more accurately the focal-conic fan-shaped texture. The ellipses are not recognizable as such. They lie in planes perpendicular to the film and are arranged along the edges of the fan-like areas. The lines following the hyperbolae appear as essentially straight lines along which the direction of extinction changes discontinuously. This change becomes smaller with the distance from the accompanying ellipse and the lines gradually disappear. Thirdly, there are the oily streaks, which are long transversely striated bands consisting of chains of focal conic groups (Fig. 2.24).152 They may appear when a well-defined focal-conic texture is disturbed by shifting the cover slip.

A homeotropic (also called pseudoisotropic) texture can also form spontaneously at the phase transition, particularly when this takes place on cooling from a homeotropic nematic. The transition is then difficult to observe except by the disappearance of Brownian motion.

64 Claudine Noel

(a)

(b)

Fig. 2.23. The separation of the SA phase in the form of batonnets from the isotropic liquid. Polymer

"'f car- CH -j- CH3 • I I

cOO-+- CH-+.2 0 -~ coo-.lQ\.- 0 -CH::;- CH-C H 2 ~ ~ _. 2 5

at (a) 14J3°C and (b) 141·g°C. (From Ref. 150.)


Fig. 2.24. Polyester

t~~ I?~~ _rr::\.J ~.1 c~o-c \:?; c-o~C-O-+CH2CH2v--t;l

displaying oily streaks and dark homeotropic regions (smectic A texture). Crossed polarizers. Original magnification x 200. (From Ref. 152.)

At the transition from a normal isotropic liquid to smectic C, batonnets may appear first similar to those of smectic A phases. They are, however, divided by inversion walls into domains of different tilt orientation. Accordingly, the smectic C batonnets show regions with different interference colours.

Smectic C modifications exhibit two microscopic textures: the Schlieren texture and the focal-conic fan texture. As previously discussed, the Schlieren texture of the smectic C phase can be distinguished from that of the nematic phase by the fact that it exhibits only singularities with S = ± 1. Compared to the corresponding texture in smectic A, the 'broken' focal-conic fan texture of a smectic C is less regular and disturbed by additional disclinations. If the C phase is formed on cooling a smectic A phase, then the Schlieren texture will be obtained from a homeotropic A texture and the broken fan texture will be obtained from the simple focal-conic texture of the preceding A phase.

Smectic C phases can be twisted by dissolving an optically active compound in them 121 ,153 or by introducing chirality into the mole-

66 Claudine Noel

cular structure.146-148,154-157 This may lead to the formation of the typical twisted smectic textures (Fig. 2.25). The striation can be explained by the periodicity of layers with equal twist angle, the separation of the stripes corresponding to the half-pitch of the smectic screw.

Transitions with the participation of liquid crystals may show characteristic phenomena. If a nematic modification turns to a smectic A or smectic C phase, transient stripes in the form of a myelinic texture (also called chevron texture or striated texture) are often visible. 95 ,116 Another example is the paramorphotic arced focal-conic fan texture of the smectic

Fig. 2.25. Twisted smectic C texture obtained by the addition of terephthalylidene-bis-4-(( + )-4'-methylhexyloxy) aniline to the Sc phase of polyester:

t[-@-@-@_~-<r-t-CH2CH2<>4J (From Ref. 121.)

E phase obtained on cooling the simple focal-conic texture of the smectic A phase. At the point of transition to the E phase, the focal-conic fans of the preceding A phase become crossed with concentric lines or arcs running across the backs of the fans (Fig. 2.26). The arcs are not transitory in nature and remain throughout the temperature range in which the phase persists.


Fig. 2.26. Paramorphotic arced focal-conic fan texture observed upon cooling from the smectic A state (cf. Fig. 2.22) for the polyester prepared from din-propyl-p-terphenyl-4,4"-carboxylate and HO~CH(CH3)~CH(CH3)~OH.

(From Ref. 149.)

As mentioned earlier, if the LC is composed of molecules having a strong longitudinal dipole moment, then it can form different types of smectic A and smectic C phases and can exhibit re-entrant behaviour. Several terminally cyano-substituted SCLCPs are presently known to exhibit the unusual 1-+ N -+ SAd -+ NR sequence.96 , IS8- 160 Typically, the high-temperature nematic phase separates from the isotropic liquid on cooling in droplets; these grow and join together to form a Schlieren texture (Fig. 2.15). Further reduction in temperature produces a chevron texture with characteristic transition bars (Fig. 2.27(a)). This texture changes on standing for some time into the stable focalconic fan texture which is consistent with a smectic A phase. As the temperature falls, either a Schlieren texture is progressively restored or a paramorphic fan-shaped texture is formed. Polymers with long spacers exhibit complex polymorphism involving different SA and Sc phases. 1S8 ,159 Usually, the transitions are difficult to observe, but in some cases subtle textural modifications such as a multiplication of the focal conics can be evidenced (Fig. 2.28).

68 Claudine Noel

(a)

(b)

Fig. 2.27. (a) High-temperature nematic-smectic A transition, T= 122-SoC; (b) smectic A phase, T = 121°C; (c) re-entrant nematic phase, T = 7S0C (c.f. following page) polymer

-+ c~ -CH--r-

I C()(}-fCH?6°-@--@--CN

(From Ref. 96.)


Fig. 2.27.-Contd.

(a)

Fig. 2.28. (a) High-temperature and (b) low-temperature (c.f. following page) smectic A phases of polymer

+cl2-flH-

tr?~ o--t- c~-c~-o-----+-@-coo---0)--c~ CHz@--cN (From Ref. 159.)

70 Claudine Noel

(b)

Fig. 2.28.-Contd.

2.5 MISCIBILITY TESTS

As mentioned earlier, texture observation cannot always be counted on to produce an unambiguous determination of the type of mesophase present. It can happen that similar textures are observed with two LC states separated by a phase transition. It is advisable, therefore, to use this method in conjunction with other types of characterization such as miscibility tests. Sackman and his school at Halle have produced a classification scheme for LMMLCs based on the criterion of complete miscibility of identical phases.74 Isomorphous LCs are considered as equivalent and characterized by the same symbol. Therefore, the mesophases of a new compound can be positively identified by isomorphy with known mesophases of reference compounds. Assuming that the method is applicable to LCPs, the type of mesophase can be determined by establishing the isobaric temperature--composition phase diagram of a binary system composed of the polymer and a reference compound. Such diagrams can be generated from thermal data or, because of the various optical features characteristic of each mesophase


structure, from microscopy observations. The latter method, called the contact method,161 allows great rapidity in the assessment of the phase diagrams.

Chronologically, the works that should be mentioned first are those of Millaud et al. 162 and Noel et aI. 121 , 163 who have determined the nature of the mesophase of a polyazomethine by establishing its isomorphy with LMM Schiff's bases of similar structure. The binary diagrams are similar to the typical phase diagrams encountered for mixtures of two LMM nematogens: an eutectic is seen in the solid-nematic curve and there is a continuous pathway from the nematic phase of reference compound to the mesophase of polyazomethine. However, only a small part of the nematic-isotropic curve could be obtained due to the decomposition of the polyazomethine well below its clearing point. Later, the applicability of the rule of selective miscibility was confirmed for several nematic MC polymersl64-167 and copolymers.97.101 By combining d.s.c. measurements with texture observations, it was possible to get diagrams of state with an uninterrupted series of mixed crystals between modifications of the same type. An example is given in Fig. 2.29. The eutectic

300

275 (,)

-:- 250 .. :; ~ 225 .. a. E 200 ~

175

o M

Isotropic

Nematic

Solid

40 80 P

Weight percent P

o 0 M-CH30% C-O-@-0-C-@-OCH3

P={6-@-O-{CH2ko-@-KO@-O(CH2~OO%Ot

Fig. 2.29. Isobaric phase diagram of a MCLC polyester and a LMM reference compound. (From Ref. 164.)

72 Claudine Noel

composition and temperature can be calculated using the Schroedervan Laar equation.16S.l69 However, the calculated values do not agree with those obtained experimentally, the discrepancy being explained as arising from difficulties in packing long macromolecular chains with small molecules. More recently, cholesteric polyesters based on 4,4'azoxybenzene and 4,4'-azoxy-2,2'-methylbenzene were found to be miscible with nematic polyesters of similar structure and an example of perfect solid, cholesteric and isotropic liquid solutions formed by polymeric enantiomers was reported. 170 Mutual miscibility tests have also been used successfully to identify SA139.l45 and Se12l phases in MCLCPs. Unlimited and limited miscibility of smectic modifications are illustrated in Fig. 2.30. It is to be noted that Krigbaum et al. 166 and Bosio et al. 145 were unsuccessful in their attempt to identify from miscibility

y 300

()

'--~ :::I ..

300

~ 200 E ~

lBBA

Fig. 2.30. Binary mixtures of a MCLC polyester and LMMLCs. (From Ref. 145.)


tests the SH phase of poly( p-biphenylsebacate) and the SE phase of poly(terphenyl-4,4/1 dicarboxylate), respectively. Thus, it appears that, for the identification of polymeric mesophases, mutual miscibility is the method of choice for nematics and smectics of low order while X-ray diffraction may be required for the more ordered smectics.

The study of binary mixtures of LMMLCs has attracted considerable attention since the earliest days of LC research. Mixtures have played an important role in the technology of LC display devices by extending the range of the nematic phase to room temperature. Moreover, extrapolation of the transition temperatures of binary systems involving a nonmesogen sometimes provides an estimate of the virtual transition temperature of the latter component. This is a quantity of theoretical interest and one which obviously cannot be determined experimentally from study of the non-mesogen as a single component.

Polymeric systems containing two different mesogenic repeating units can be prepared in various ways. A mechanical mixture of two meso genic homopolymers is the direct analogue of the mixtures of LMMLCs. However, the two types of repeating unit can be mixed more intimately within a single chain by random copolymerization, and still other possibilities are block and graft copolymers. Krigbaum et aI.171173 have investigated the mesomorphic properties of binary systems formed by polyester mixtures and random copolyesters in the hope that these might provide additional insight into the nature of polymeric mesophases at the molecular level. A prominent feature of their results is that cocrystallization is fairly frequent for LC copolymers. However, copolymerization disrupts the structural regularity of the crystal, decreases the degree of crystallinity and broadens the crystal melting transition. These effects become more pronounced with increasing differences between the two repeating units. Moreover, copolymerization lowers the crystalmesophase transition temperature, which increases the thermal stability range of the meso phase. Both nematic and smectic phases can accommodate different units of the same mesogenic type with minimal disruption of the mesophase structure. The variation of the nematic (or smectic) - isotropic transition temperature and the corresponding entropy and enthalpy changes very nearly follow the arithmetic averages over those of the parent homopolymers. These conclusions agree with similar experimental data reported from several other laboratories. 129,174-180 It demonstrates that copolymerization not only represents one of the most powerful techniques available for manipulating polymer phase transitions but that it also becomes an important ex peri-

74 Claudine Noel

mental tool that can be used to obtain information about eventual virtual transitions of non-meso genic homopolymers.175.178.179

The applicability of the rule of selective miscibility has also been examined for mixtures of SCLCPs with conventional nematogens.181-187 Unlimited miscibility is illustrated in Fig. 2.31. If the polymer is essentially non-crystalline in character, no eutectic is seen: the melting point of the reference compound is depressed as the polymer concentration is increased, whereas the glass transition of the polymer is lowered as the mole fraction of the LMMLC is increased, this being the so-called plasticizing effect. As a consequence, at certain compositions the thermal stability range of the mesophase is noticeably enlarged towards lower temperatures. However, if the polymer is semi-crystalline, the melting points of both components are depressed and a eutectic is seen, indicating the presence of distinct crystal structures for the two components.

Re-entrant behaviour188 and enhanced or induced SA phasesl89.190 have been observed in some mixtures involving a LMMLC and a SCLCP. An important point to be noted here is that the components were non-polar. Needless to say, we are nowhere near explaining these facts from a molecular-theoretical point of view. From the phenomenological point of view, these phenomena can be described as arising from anomalies of periodicity. On the basis of X-ray experiments it has been suggested that these SCLCPs exhibit partially bilayer SA phases191 and that changes in the anti parallel arrangement of the side chains are responsible for these effects.

In contrast to LMMLCs and MCLCPs, unlimited miscibility seems to exist only if the chemical structure of the polymer-bound side groups is similar to that of the LMM reference compound. The first observation of a heterogeneous region between two nematics was made in 1982 by Casagrande et al. 185 In 1984, Brochard et al. 192 combined the Maier-Saupe theory of the nematic-isotropic transition and the FloryHuggins theory of mixtures to describe the phase diagrams of two nematogens A + B. For certain sets of parameters, rather complex diagrams were predicted when one component is a SCLCP. The miscibility in the nematic state is very small as soon as the nematic interaction parameter U AB is smaller than U AA (or U BB). This is related to the most familiar case of a conventional polymer in a nematic solvent. 193

On the other hand, if U AB > U AA, the miscibility of a SCLCP and a LMMLC is largely increased by the nematic alignment. Reliable data on the effects of change in molecular structure on the SCLCP/LMMLC phase diagrams confirm the conclusions reach about the role of AI A, BIB

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76 Claudine Noel

and AlB interactions. 186,187 More recently, a density-functional expansion method was used to calculate the phase diagrams of LCPs. 194,195

Special attention was paid to the influence of a number of parameters (molecular weight, persistence length, isotropic and anisotropic interactions). This approach appears quite general and applicable to all types of LC chains simply by appropriate definition of the correlation functions.

Copolymers known as partially fixed SCPs are of great help in overcoming the drawbacks encountered with fully substituted chains. An interesting example is the partial substitution on a poly(methylsiloxane) backbone which strongly promotes the solubility in a given LMMLC compared with cases of totally substituted or unsubstituted chains. 196

Such copolymers also have the advantage of lower viscosities. 197,198

Hardouin and colleagues have undertaken an investigation of phase diagrams of binary mixtures involving a partially fixed SCLCP and a LMMLC in the hope that these might provide additional insight into the nature of the SA polymorphism in high-molar-mass systems. For recent results see Refs 197 and 198.

2.6 X-RAY DIFFRACTION PATTERNS

The fascination and the difficulty of structural studies of the mesophases is that elements of crystal symmetry exist together with considerable disorder characteristic of liquids. This means that no simple formalism is appropriate for all phases, and the difficulties are further compounded because the molecules are large and non-rigid. Hence, most of the diffraction work reported for LCPs has been aimed first at characterization or verification of the phase behaviour, where it is extremely valuable and generally unambiguous, and then determination of layer and molecular spacings and correlation lengths, which can give valuable insight into the molecular packings and a measure of the extent of order. These studies require, in essence, only measurement of positions and widths of the diffraction features. A small number of more detailed investigations have also been made to try to determine quantitative aspects of the distribution functions. These studies require the measurement of the intensity of scattering in reciprocal space and are consequently much more difficult and time consuming. This subject has been reviewed in a number of reviews more or less specific to MC50,149,199-201

or SC55,126,150,202-211 LCPs and I shall concentrate attention on the more disordered phases: nematic, smectic A and smectic C.


2.6.1 X-ray Diffraction Patterns for Powder Samples For powder samples, the well-known Debye-Scherrer technique is used. This method gives all the reticular spacings but no information about the spacial orientation of these planes. The two main diffraction effects usually observed on the X-ray diffraction pattern of a powder sample are the 'inner ring(s)' (close to the centre of the diffraction pattern, and corresponding to a diffraction angle e of only a few degrees) and the 'outer ring(s)' (much farther out in the diffraction pattern and corresponding to a diffraction angle e of about 10°). The inner rings are indicative of longer layer spacings. The outer rings correspond to shorter preferred spacings occurring in the lateral packing arrangement of the molecules. The appearance of a broad halo or a sharp ring furnishes a qualitative indication of the degree of order.

X-ray studies on thermotropic LMMLCS72•73 ,212.213 have shown that the mesophases can be divided into three groups according to the characteristics of their X-ray patterns at large diffraction angles. In the first group we find the N, SA and Sc phases whose diffraction patterns show only one broad, diffuse halo indicating a lack of periodic lateral order: the distribution of the molecular centres of mass is random. The second group is composed of the 'crystalline' smectics (SB, SE> SG, SH, SJ and SK)' Their diffraction patterns show a single or several sharp outer rings which are related to the high degree of order within the layers. If the molecules form a hexagonal close-packing of cylinders as in the SB phase, the distances between neighbouring molecules are equal, and there is only one sharp diffraction maximum corresponding to the nearest-neighbour distances. If there is more than one such maximum, then there must be also more than one value for these distances, and the packing can no longer be hexagonal. Thus, the number of maxima gives information about the type of the molecular packing. The SBHm SF and Sf phases are intermediate between these two groups.

The X-ray patterns of nematics and smectics differ mainly in their characteristics at small diffraction angles. Nematic patterns present a diffuse ring, indicating that there is no order in the direction of the molecular long axes. It should be noted, however, that in the nematic phase the inner ring is stronger and less diffuse than in the isotropic phase. It corresponds to distances which are generally quite close to L, the length of the molecule calculated from standard bond lengths and angles under the assumption of a planar, all-trans conformation of the molecule. In contrast, smectic patterns present one or several sharp rings, showing the existence of layer-like correlations. These diffraction maxima

78 Claudine Noel

provide direct information about the layer thickness. For orthogonal phases, i.e. phases in which the director is perpendicular to the smectic layer, one might expect that the layer spacing, d, would be equal to the molecular length, L. It has generally been found, however, that d is less than L, and various explanations have been proposed for this.69,213 It is worth noting that the differences between d and L can be explained by taking account of the orientational disorder of the molecular long axis which is present in all smectic phases. For 'tilted' phases, the differences between d and L may be used to calculate the 'tilt' angle. However, there is the possibility of some ambiguity because of unknown molecular conformation.

For both MC97 ,109,119,131,214,215 and SC 55 ,95,143,204,206 LCPs, the

literature reports X-ray diffraction patterns which are compatible with a nematic structure. They present at large angles a broad diffuse halo arising from the intermolecular spacings perpendicular to the long axes of the meso genic groups. The average intermolecular spacings obtained by the rather arbitrary use of the formula d = 1·2),./2 sin e based on arguments of cylindrical symmetry in all cases lie in the range 3-5-6'5 A. Such values correspond approximately to the average widths of the molecules but are certainly smaller than the diameter of a freely rotating molecule.

In 1986, Duran et al.216 reported for a LC polymethacrylate an X-ray diffraction pattern different from those obtained for conventional nerna tics. It was characterized by two diffuse rings corresponding to d spacings of about 30 A and 4·7 A, respectively, and a sharp reflection giving a d spacing of 8·4 A. The diffuse ring at small angles was interpreted as an indication of cybotactic ordering; it corresponded to twice the extended model length of the side chains. The diffuse ring at large angles was associated with the lateral spacing between the mesogenic side groups. This phase was designated as a 'novel nematic' meso phase, though its detailed structure was left unresolved, the origin of the sharp diffraction at 8'4A being obscure. In 1987, Duran et al. 217

proposed the concept of ribbon-like conformation of polymer chains and it became clear that the sharp reflection at 8·4 A represents a stacking periodicity related to the thickness of the ribbons (Fig. 2.32). It is worth pointing out that this structure is intermediate between smectic and nematic or columnar, as it corresponds to the periodical stacking of two-dimensional palisades or ribbons.

SA and Sc meso phases have also been identified in several MC 144,145,153,214,215,218-220 and SC 55 , 126,143,146,147,155,156,191,204,

221-225 LCPs using X-ray diffraction methods. Typical diffraction


o

30 A +-- ------ •

Fig. 2.32. Structure of the 'novel nematic'. The darkened part of the ribbons corresponds to the polymer backbone.

(From Ref. 217.)

patterns of powder samples are shown in Fig. 2.33. They consist of a diffuse outer ring associated with the unstructured nature of the layers and a sharp inner ring (with sometimes its second order) related to the lamellar thickness. Diffraction patterns of unoriented Sc phases are essentially the same as those for SA except that the layer spacings, d, are relatively smaller. Indeed, the accepted structure for the Sc is that the director is tilted to the layers while in an SA phase the director is normal to the layers. Usually, if the smectic C phase is followed by an A phase the tilt angle decreases gradually and finally becomes 0-100 at the C-A transition. 146, 147,191

It should be noted that a number of SCLCPs are now known where d> L. Polymers with terminally cyano-substituted chains are particularly susceptible to this behaviour.96,143,191,198,204,222,224,226,227 From simple

energy considerations it is evident that interactions between neighbouring dipoles can by no means be neglected in such strongly polar materials. The highly polar cyano group attached at the end of the side chains results in strong anti parallel near-neighbour correlations. As a consequence, the smectic A phases of these polymers often consist of 'bilayers', the molecules arranged in an anti parallel overlapping interdigitated structure.

Similar, if less pronounced, effects are also found in other non-polar SCLCPS. 191 ,223 A simple estimate of dipolar-induced dipolar forces in these systems suggests that the core overlap is due to the action of these forces between the meso genic side groupS.228 SCLCPs may exhibit A and C phases which have a monolayer, a perfect bilayer or an interdigitated structure, depending on the nature of the polymer backbone, the number of atoms in the flexible spacer and the length of the terminal group.146, 147,205,206,229

A remarkable dependence of the layer spacings, d, on the degree of substitution was reported for different poly(methylsiloxaneco-dimethylsiloxane)s containing non_polar230-232 side groups. For the

80 Claudine Noel

(a) (b)

(c)

Fig. 2.33. Typical X-ray diffraction patterns for unoriented (a) nematic, (b) smectic A, and (c) smectic C phases.

homo polymers in the smectic A state, the layer spacings correspond approximately to the lengths of the side groups calculated with the assumption of an all-trans molecular conformation, which is consistent with the formation of monolayer structures. When the concentration of the structural units containing meso genic groups decreases, the layer thickness increases. This dependence of the lamellar thickness on copolymer composition was explained by a microphase-separated copolymer morphology. Such a morphology requires a distortion of the


random coil conformation of the flexible backbone to the extent that it can be squeezed in between the smectic layers. Quite recently, similar effects were reported for poly(methylsiloxane-co-dimethylsiloxane)s containing polar side groupS.198 However, the increase in d was significantly less than that observed in the experiments described above for non-polar systems. This could be understood by considering a SAd structure with d > L for the homopolymers.

More ordered smectic phases have also been identified in several MCI45.166.233.234 and SC55.191.206.217 •. 235-240 LCPs. The paper by

Freidzon et al. 237 is one of the most recent examples of such studies. The X-ray diffraction patterns obtained at different temperatures for the SCLC polymer

---t CH2-CH4-I

o Ao+ CH2--l-S coo--@-coo-@-- OC3 H 7

(annealing at S 60°C S 95°C N 116°C I 350C for several --+ \G I C ---+ --->

days) /400C SF

reveal three smectic phases in addition to a nematic phase. The Sc phase with unstructured layers gives a diffuse maximum at large diffraction angles. The intralayer correlation length, (.1, calculated from the halfwidth of the maximum by assuming a Lorentzian line shape is of the order 5-10 A. The SF phase gives a maximum which is about half as broad as in the Sc phase. This is consistent with the higher degree of order within the layers of 40A. For the more ordered SG phase a sharper maximum and additional weaker reflections are observed. The intralayer correlation length is approximately 70 A, which is smaller than that determined for LMM smectics G. This indicates noticeable disorder of the molecules.

Quite recently, smectic E phases were identified for a polymethacrylate and two alternating copolymers of maleic anhydride with vinyl ethers, bearing a 4'-methoxy-4-biphenylyl group via an oligo (ethylene oxide) spacer. 217 •. 238 Smectic layers are single layers of ribbon-like polymer chains with all pendant groups arranged in a single row and arranged on the same side of the backbone. Mesogenic groups are oriented up and down at random and polymer chains are aligned along the [110] direction of the rectangular lattice describing the packing of the pendant groups.

82 Claudine Noel

2.6.2 X-ray Diffraction Patterns for Oriented Samples If a sample can be obtained in the form of an oriented monodomain, it is possible to extract more detailed structural information from its diffraction diagram. 72 •73 For nematics, monodomains can be obtained by orientating a powder sample in a strong magnetic field. Aligned SA and Sc may then be prepared by careful cooling from the aligned nematic phase. An alternative procedure, useful for preparing monodomains of the more ordered phases, is by drawing fibre out of the meso phase and quenching the LC array in the glassy state so that it can be examined at room temperature.

Examples of oriented nematics can be given for both MC97 ,101,129,214d,241-243 and SC 54,55,96,206,209,211,244-250 LCPs. A typi-

cal diffraction pattern is given in Fig. 2.34. The anisotropy is clearly shown and there are correlations of two distinct periods perpendicular and parallel to the director n. For Q 1- n the diffraction pattern is very liquid-like. The outer diffuse halo evidenced for powder samples is split into two symmetric crescents. Their angular extension reflects the degree

(a)

(b)

Fig. 2.34. Typical X-ray diffraction patterns for oriented samples: (a) conventional nematic, and (b) nematic with cybotactic groups.


of parallel alignment. The dominant features of the scattering with Q II n are arcs or short bars which are associated with intramolecular interferences. They appear distinctly only when films are overexposed, since their intensity is smaller than that of the strong crescents.

If the samples are aligned by cooling in a strong magnetic field from the isotropic liquid into the nematic phase, then the two crescents are located in a direction perpendicular to the field, indicating that the meso genic groups are more or less oriented along the magnetic field. In the same way, if MCLCPs are stretched in the nematic phase, the meso genic groups are parallel to the fibre axis, the macromolecular chains being preferentially aligned in the direction of extension. The situation is more complex for SCLCPs. Results obtained so far suggest that the relative orientation of the side chains and the polymer backbone with respect to the fibre axis depends on the specific chemical structure. In most cases, the meso genic groups are parallel to the fibre axis. The backbones probably have a random conformation. However, in few X-ray patterns the two crescents are located in a direction parallel to the fibre axis, which is consistent with the side chains perpendicular to the stretching direction.244

Secondary nematic structure was found by Blumstein and coworkers129.243 during the X-ray investigation of polyester prepared from 4,4' -hydroxy(2,2' -methyl)azoxybenzene and dodecanedioic acid. The diffraction pattern of the nematic shows the development of enhanced order characteristic of smectic C phase: the first of the meridional arcs splits up into four spots. This phenomenon is incompatible with the classical definition of the nematic phase and suggests an additional order of macromolecules within cybotactic groups for which a structural model was proposed. Cybotactic nematic mesophases were also found for the copolyesters of the same series containing spacers with an even number of methylene units, but ordinary nematic phases were identified for the members with odd-numbered spacers.25i

Noel et al.97 used X-ray diffraction to characterize the nematic phase of a co polyester based on terephthalic acid, methylhydroquinone and pyrocatechol. Along the equator, in addition to the typical crescents, the X-ray diffraction patterns showed two diffuse spots at smaller angles. Similar diffraction features have been reported previously for certain LMM nematic systems, but without any justification.2s2 Taking into account that such diffuse spots have been observed for helical structures,253.254 Noel et ai., in their analysis of the X-ray patterns, considered the possibility of such arrangements of the chains but without long-range

84 Claudine Noel

order. They proposed a structural model which bears a resemblance to the molecular arrays accepted for cybotactic phases. The two diffuse spots would be expected for roughly parallel chain bundles in the form of two/four-stranded ropes.

Cybotactic nematic phases were also reported for SCLCPS. 55 ,244, 246-248,255 These are polymers in which pronounced smectic order parameter fluctuations occur. This is especially true of polymers which exhibit an SA or a Sc phase at lower temperatures.

As already mentioned, a number of LCPs adopt layered structures in the melt but, compared with their LMM analogues, LCPs exhibit mainly modifications with unstructured layers, i.e. SA and/or Sc. Typical X-ray diffraction patterns for oriented samples of an SA and an Sc phase are shown in Fig. 2.35. As the structure within the layers is disordered, the symmetrical crescents at large angles are qualitatively those for a nematic. However, very strong sharp reflections are seen at small angles instead of diffuse arcs. These Bragg spots show the existence of extensive layer-like correlations.

Different types of ordering may be obtained depending on the materials and the conditions. If the samples are oriented in a magnetic field, then the relative position of the large-angle crescents and the small-angle Bragg spots with respect to the direction of the field is indicative of mesogenic groups parallel to the field with perpendicular (SA)230,247.249.256-259 or tilted (Sc)248 layers. On the other hand, when MCLCPs are stretched in the smectic A or smectic C state, the macromolecular chains tend usually to align in the direction of

(a) (b)

Fig. 2.35. Typical X-ray diffraction patterns for oriented (a) SA, and (b) Sc.


extension. As a consequence, the smectic layers are perpendicular to the fibre axis and the mesogenic groups are arranged with their long axes perpendicular (SA)153 or tilted at a noticeable angle (Sd 141 ,153 to the layer planes. However, parallel orientation of smectic layers and fibre axis was reported in a few cases.220 The situation is again more complex for SCLCPs. In most cases, when the samples are mechanically aligned by stretching the polymers, the layers are parallel to the direction of extension. 146,147,209,244,245,260 However, the results reported so far indicate that the type of ordering is dependent on the specific chemical structure, the degree of polymerization261 and the nature of the meso phase from which the fibre is drawn.245 For the former, an influence of the length of the flexible spacers was recently reported by Davidson and Strzelecki256 for the polyacrylates of formula:

-4 fH- CH~ rO-+CH~ o-~-0--0-C-0-CN

2n II~ II~ o 0

The X-ray diffraction patterns obtained for the members with long flexible spacers (n = 8, 12) were consistent with the layers parallel to the stretching direction and the mesogenic groups perpendicular to the fibre axis. The opposite effect was observed for a poly acrylate of the same series, but with shorter flexible spacers (n = 6). In this connection it is worth noting that SCLCPs having meso genic groups with ring ~ CN terminal substituents and short flexible spacers in general align under stretching with the side chains parallel to the fibre axis. 247 ,255 This is probably related to the anti parallel, overlapping interdigitated SAd structure of these strongly polar materials.

In addition to the diffuse crescents and the Bragg spots, the X-ray patterns for oriented SA and Sc SCLCPs can show parallel diffuse lines equally spaced versus Q.147,247,248,260 This is characteristic of disorder along the director. The extra scattered intensity arises from uncorrelated periodic columns which are out of the mean positions in the layer plane. 73 Such diffuse features do not usually appear in SA and Sc diffraction patterns of LMMLCs. They often occur in smectic phases which exhibit three-dimensional order, and so the most obvious explanation of their appearance in the X-ray diffraction patterns of SCLCPs is in

86 Claudine Noel

terms of rigidity effects of the smectic layers due to the macromolecular nature of the compounds.

A second diffuse zone can be seen:246.247,262 four off-meridian spots are visible. Konstantinov et al. 246 assigned these reflections to the formation of 'blocks' of meso genic groups. By comparison, this extra scattering was interpreted by Davidson et al. 24 7 to be caused by a periodic modulation of adjacent layers, with a wave vector parallel to the layer plane.

Different models of the SA and the Sc phases stemmed from the numerous investigations on SCLCPS.55,206,210,230,245,247,263 In

general, the same arrangements as known from LMMLCs were proposed for the packing of the mesogenic side chains without making any definite proposals about the conformation of the polymer backbone. Recent X-ray diffraction studies of polysiloxanes257 ,264 polyacryiates240 and polymethacrylates239,24o showed that the smectic layer is divided into sublayers of different electron densities consisting of the meso genic cores, the spacers and the terminal substituents, respectively; the polymer backbones are confined between the smectic layers in two dimensions. However, values determined by small scattering methods (see Section 2.7) for the radii of gyration of the main chain in directions parallel and perpendicular to the director show that the main chain cannot stay strictly confined between two layers. This was theoretically treated 265 by assuming a hopping process of the backbone from a layer to an adjacent one, thus creating defects in the smectic planes. Quite recently, detailed X-ray diffraction studies by Davidson and Levelut266 have conclusively proved the existence of such defects.

It should be noted that Zugenmaier and Miigge267 deduced a different model from X-ray investigations of polysiloxanes with mesogenic side groups in the crystalline and smectic state. According to these authors, the polymer backbone forms an ordered structure from which the mesogenic groups stick out at an angle of approximately 90°, depending on packing effects and chemical constitution of the side chains. The macromolecules (backbone and side chains) form elliptical bodies which can be shifted with respect to their long axes and which are not or only weakly correlated. Although this model, which assumes a well-ordered conformation for the polymer backbone, does not seem to agree with recent small-angle neutron scattering studies (see Section 2.7), it may be considered valid for some polymers.

Characterization of Mesophases

2.7 CONFORMATION OF LCPs AS REVEALED BY SMALL-ANGLE SCATTERING METHODS

87

Polymer entropy is antagonistic to nematic (or smectic) order. A flexible polymer is expelled from a nematic solvent, since a chain stretched out to be parallel to a nematic field loses a lot of entropy; equally the nematic field, distorting to be locally parallel to a random polymer, would develop a large nematic elastic energy.

In the case of MCLCPs, depending on the size of the spacers, we may go from very rigid to very flexible chains. But it is not only the chain that matters. As mentioned above, the surroundings are also important. In 1981 de Gennes considered the possibility of significant conformational changes at the isotropic-nematic transition. A spontaneous change in the chain shape at the transition has been predicted, leading to an increase in the chain dimension parallel to the average orientation and a decrease in the perpendicular direction.268 ,269

A variety of experimental techniques have been used to measure the orientational order in MCLCPs. A prominent feature of the results is that the orientational order of the spacers is considerably larger than that of chains appended to LMMLCs. The parallel alignment of rigid core parts of the repeating units constrains movement within the spacers to some extent and, although the spacers are not restricted to the all-trans conformation,270-273 the conformations that are adopted are rather extended.270,271,274276 The efficacy of such enhanced ordering in MCLCPs is dependent on the molecular characteristics of the spacers (their length, their flexibility, ... ) and the nature of the functional groups which link the spacer segments to the mesogenic core. In particular, the population of highly extended conformers is strongly correlated with the number of bonds--even or odd-in the spacers. Similar conclusions were drawn from conformational energy calculations.277 280 The paper by Yo on and Bri.ickner278 is one of the most recent examples of such studies. By examining the distribution of chain sequence extension for three types of MCLCPs, these authors found that characteristics of chain sequence extension and extended conformers relate very closely to the experimental results of isotropic-nematic transitions of the polymers. They demonstrated that detailed conformational order of flexible segments in a nematic state can be deduced by matching the experimental results of enthalpy and entropy changes with those estimated from the conformational selection on the basis of chain sequence extension.

88 Claudine Noel

Satisfactory methods for the measurement of the degree of chain extension are still being explored. Reliable data are available for only two nematic Me polymers:

where n=7 and 10.281 ,282 In agreement with theoretical predictions,268,269 small-angle neutron scattering (SANS) experiments showed that in the nematic state these polymers might be considered as coils having the symmetry of flattened rotational ellipsoids with the long principal axis in the direction of the director. Both the radius of gyration, RG , in the isotropic state and the radius of gyration parallel to the director, R 1i , in the nematic state increase as the molecular weight is increased.

The polymer backbone of a SCLCP has a strong requirement, if it has any flexibility at all, to be as random as possible in its configuration of shape. This conflicts with any requirement of orientational order that nematogenic groups, present as attached side chains, might have. Kunchenko and Svetogorsky283 and Wang and Warner284 modelled this competition taking main chains of various stiffness and side groups of various length. The antagonistic influences of the nematic field and the entropy of the chain were resolved by a distorsion of chain statistics away from spherical. The theory predicts molecular conformational changes at the isotropic-nematic transition. Depending on temperature, nematic coupling and stiffness, the chain becomes prolate (rod-like) or oblate (disc-like) (Fig. 2.36).

Fig. 2.36. Three possible phases for comb nematics. (From Ref. 284b.)


A statistical average of the clearing entropies of over 100 LCS285 ,286 yielded values of approximately 2 ± 1, 7 ± 5 and 15 ± 7 J K -1 mol- 1 for LMMLCs, SCLCPs and MCLCPs, respectively, This is related to the degree of extension of the flexible segments, As already discussed, in MCLCPs flexible spacers do not play the role of solvent but rather participate actively in the ordering process, As a consequence, they exhibit rather extended conformations. Less effect is observed for SCLCPs. Indeed, even though the spacers experience on one end the ordering tendency of the mesogen, they are linked at the other end to a more or less flexible macromolecular chain that tries to achieve a random coil conformation. From a detailed analysis of deuteron NMR spectra, Spiess et al. 287 289 established that the order parameter, S, decreases from mesogen to spacer to polymer backbone. However, local geometric factors have to be known before one can translate these measurements into a gross order parameter characteristic of the oblate or prolate spheroidal shape of the chain as a whole.

Experimental studies of the magnetic field effects on the SCLCPs

m=3-6

either in the melt290 or dissolved in a LMM nematic99 ,291.292 were recently reported. While the static properties (the elastic constants K 1

and K 3 ) were similar to those of LMMLCs, the hydrodynamic behaviour resembled that of conventional polymers in the melt. The analysis of the data, within the framework of a hydrodynamic model due to Brochard,293 indicated that the polymer backbone would have a nonspherical conformation, the anisotropy defined by the ratio R:c/RI of gyration radii being large enough to be measured by a small-angle scattering method. A comb polymer in a nematic solvent can offer sufficient contrast that its shape can be seen by small-angle X-ray scattering (SAXS), obviating the need for deuteration. Although it is not a probe of pure SCLCP behaviour, Mattoussi et aJ.294 exploited the observed difference between the mean sizes of the polymer perpendicular and parallel to the nematic ordering direction to deduce the nematic coupling to the main chain. The parallel dimension was greater than the perpendicular one. For a polysiloxane of similar structure (the -COO- groups are reversed to -OOC-) but in the melt, Moussa et al. 295 determined by SANS a ratio R:c/RI; of gyration radii of 0'75, in close agreement with the value of 0·73 reported by Mattoussi et al.294

90 Claudine Noel

These results might serve as an experimental corroboration of the prolate conformation predicted by Wang and Warner284a for the NIl or NIll phase (Fig. 2.36).

Of special interest is the fact that, in the smectic state, the anisotropy depends on the degree of polymerization.296 For short chains, R~ is smaller than R!I in both the nematic and the smectic A phases. By contrast, for long chains, the anisotropy reverses itself at the nematicsmectic transition so that the polymer backbone seems to adopt an oblate conformation in the smectic state. It should be noted, however, that the conclusions drawn from these SANS experiments must be considered with great caution, the polysiloxanes under interest having been labelled in the terminal methoxy group of the side chains.

Boeffel et at. 297 ,298 inferred from recent 2H-NMR work on a polyacrylate and different polymethacrylates that the local orientation of the polymer backbone depends on the specific chemical structure of the backbone. Their results on the poly acrylate are consistent with the NIll phase predicted by Warner et at.: a parallel orientation of the polymer chain and the side chains is preferred. In contrast, for the polymethacrylates, the perpendicular alignment of the backbone and the side chains is consistent with the NJ, phase. It should, however, be pointed out that recent SANS experiments299.300 suggest that the conclusions derived for the polyacrylate by Boeffel et ai. are misleading. As already mentioned, NMR yields information about local orientation and the complexity of the systems investigated makes it difficult to interpret the data in terms of overall conformation of the main chain.

Several different polyacrylates299 ,300 and polymethacrylates 295 ,296, 301-304 have also been studied by SANS. Experiments so far point to a perpendicular coupling (R~ > R II , oblate shape) in the nematic state and in particular to the N1 phase predicted by Wang and Warner.215 It should be noted, however, that the degree of anisotropy is rather small. This suggests that the mutual coupling between the backbone and the mesogenic side groups is only weakly negative for these systems. Quite recently, Kunchenko and Svetogorski305 using a simple model (a comblike freely jointed chain) have made an estimate for the second moment of the distribution of the meso genic units relative to the backbone segments. For the experimentally measured values a= R~/RII = 1'1 302 or 1.25301 and an estimate for the nematic order parameter S" of 0'5, the parameters, Sf, reduced to the minimum of -1/2, are 1/3 and 1/2, respectively.

In the smectic state, the degree of anisotropy is substantially greater than that in the nematic phase. The backbone appears more or less


confined in one or two layers. Both radii of gyration are temperature dependent. The behaviour of R.l differs noticeably from one SCLCP to another. By contrast, in most cases, the variation of RI' can be approximated by an Arrhenius law. Good agreement is obtained with the 'layer hopping' model of Renz and Warner265 based on De Gennes' calculation:306

Rfrx exp( -E/kBT)

The activation energy deduced from this model is '" 80 kJ mol-I, which is ten times greater than that estimated for a smectic defect by Kunchenko and Svetogorsky.283

Many experimental problems continue to be of vital importance in establishing the validity or otherwise of theoretical models. Neutron investigation of chain shape in different nematic comb melts would help establish the relation, if any, between variation in hinge structure and nematic main chain and side chain coupling and the resultant phase symmetry. With the advent of a wider range of syntheses, many more proposals and speculations for new types of behaviour and effects can be made.

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Chapter 3

N M R Studies of Thermotropic Polymers

Fran~oise Laupretre Laboratoire de Physico-chimie Structurale et Macromoleculaire associe au

C.N.R.S., 10 rue Vauquelin, 75231 Paris Cedex 05, France

3.1 INTRODUCTION

As already emphasized by some review articles, among the various techniques that are used to study polymers, NMR spectroscopy has proved to be of particular interest in the case of thermotropic polymers. 1,2 The tensorial interactions such as the chemical shift anisotropy, homonuclear couplings in the case of 1 H NMR, heteronuc1ear dipolar couplings in the case of 13C NMR, and quadrupolar couplings in the case of spins higher than 1/2 such as in 2H nuclei, are averaged to zero by the rapid Brownian motions of molecules in solution. However, they are not or are only partly averaged in a polymer below its glass transition or in a meso phase. Importantly, all the physical interactions listed above are orientation dependent. They may also be partly averaged by molecular motions. Therefore, measurement of their intensities provides information which can be used in the investigation of orientational phenomena, molecular dynamics and organization in the solid state. Differences in intensity of the physical interactions also result in differences in relaxation times which can be used to perform discriminatory experiments, leading to the selective observation of nuclei having some well-defined characteristics.

Application of these concepts to mesomorphic polymer systems will be developed in the following sections. Section 3.1 will be devoted to

103

104 Franroise Laupretre

orientational and conformational NMR studies of liquid-crystalline polymers. In Section 3.2, results obtained on local dynamics of mesomorphic polymers using NMR techniques will be reviewed. In Section 3.3 examples of recent NMR investigations of collective motions will be considered.

3.2 NMR INVESTIGATION OF ORIENTATIONAL AND CONFORMATIONAL PHENOMENA IN MESOMORPHIC

POLYMERS

The molecular organization in the different meso phases has been reviewed in the first chapter of this book. In a very general way, the degree of alignment of a molecule in a given mesophase can be described in terms of an orientational distribution function which provides the probability of finding a molecule in a solid angle ranging from Q to Q + c5Q with respect to a reference frame in the sample. A detailed description of this formalism can be found in Refs. 2 5. For a qualitative understanding of the results summarized in this chapter, which mostly deals with uniaxial phases, we will confine ourselves to the simplified concept of orientational order parameters.

3.2.1 Orientational Order Parameters In nematic and smectic A and B uniaxial phases, the average direction of orientation of the long axes of the molecules defines the director D. If the liquid crystal molecules are taken as rod-like, the degree of parallel order of the individual molecules is described by a single orientational parameter:

(3.1)

where () is the angle between the individual molecular long axis and the director, and the angular brackets indicate a thermal average. As noted already in Section 1.5, the names order parameter or anisotropy factor are also used for this quantity.

For rigid non-cylindrical, elongated molecules, the order is characterized by two parameters: the first-order parameter Szz pictures the average orientation of the molecular long axis with respect to the main director; the second-order parameter, c5 = S xx - Syy, reflects the difference in ordering of the two short axes. The quantity c5/Szz is the anisotropy parameter of the molecular orientational order.

N M R Studies of Thermotropic Polymers 105

3.2.2 Principles of the NMR Measurements A complete introduction to the NMR phenomena described in this chapter can be found in Ref. 6. In the following, we will focus on the measurements that have proved most useful in the case of thermotropic polymer studies.

3.2.2.1 lH NMR For a pair of protons Hl and H2 separated by a distance rHH, the splitting flv between the two spectral lines as a result of dipolar coupling between the nuclear magnetic moments can be written as

fl v = (3'Y~h/2nr~H)< 3 cos 2 f3 - 1)/2 (3.2)

where f3 is the angle between the H 1 - H2 vector and the static magnetic field, Ho, of the NMR spectrometer, 'YH is the proton gyromagnetic ratio and h is the Planck constant divided by 2n. Therefore, measurement of flv yields the orientational order parameter associated with the Hl -H2 vector, SHI -H2' In systems with transverse isotropy, this order parameter can be related to the molecular order parameter if the conformation of that molecular site is known. A difficulty with 1 H NMR is that usually several H-H internuclear vectors are present. When their contribution can be separated, for example by using selectively deuterated molecules, discussion of the results obtained from each vector can be made in terms of both orientational and conformational order.

3.2.2.2 2H NMR Deuterium, whose spin is equal to 1, is a quadrupolar nucleus. Its NMR parameters are almost exclusively governed by the quadrupolar interaction with the electric field tensor (FGT) at the deuteron site. Although the intermolecular contributions to the electric field gradient can be significant, 7 the field gradient usually originates mainly from the electrons in the C-2H bond and is considered to a first approximation to be intramolecular. It is found to be axially symmetric about the C-2H bond in aliphatic compounds, and to a good approximation in aromatic compounds as well. Thus, information obtained from 2H NMR on molecular order and dynamics mostly concerns individual C-2H bond directions.

In the absence of motions, the NMR spectrum is a typical quadrupolar doublet, with a symmetric splitting flv:

(3.3)


where e2qQ/h is the quadrupolar coupling constant, 11 is the asymmetry parameter, and the orientation of the magnetic field in the principal axis system of the FGT is specified by the polar angles J and cpo In the simplest case, with 11 ~ 0, J is the angle between the C-2H bond direction and the external magnetic field. The magnitude of ~v is then directly related to the orientation of the C-2H bonds in the sample with respect to the static magnetic field. Therefore, it can be used to determine the order parameter of rigid, ordered systems. Moreover, for mobile ordered systems, where there is a rapid uniaxial motion of the C-2H bond around the molecular axis, the value of the splitting yields the order parameter corresponding to the C-2H bond. When the 2H NMR spectrum is well-resolved, the order parameters S of all the magnetically inequivalent C- 2H bonds can be determined. These S values provide information about the conformational order of the different C-2H groups of the molecule in the mesomorphic state.

3.2.2.3 13e NMR The major obstacle in high-resolution solid-state 13C NMR is the dipolar coupling of the carbon nuclei with neighbouring protons. When the heteronuclear dipolar broadening is removed by irradiation of the protons with a strong radiofrequency field in the neighborhood of their Larmor frequency,S the position of the 13C NMR lines is governed by the chemical shift phenomenon. The chemical shift is a tensor. Its value depends on the orientation of the electronic distribution about the nucleus with respect to the external magnetic field. For a given orientation, the observed chemical shift corresponds to the (Jzz component of the chemical shift tensor along the external field:

3

(Jzz= I (JiCOS 2 Yi i= 1

(3.4)

where (Ji are the principal tensor components and Yi the angle of the ith component with the external field. When the signals of the different 13C

nuclei of the molecule studied are resolved, values of order parameters associated with the different carbon atoms in an oriented material can be derived from each chemical shift and related to both orientational and conformational order.

High-resolution solid-state 13C NMR 9 is also a very interesting tool for studying non-oriented polymer liquid crystals. Proton dipolar decoupling (DD), which suppresses the 13C_1 H dipolar broadening, and rapid magic-angle spinning (MAS), which averages the chemical shift anisotropy, allow the recovery of high resolution and are of major interest for the

N MR Studies of Thermotropic Polymers 107

determination of the chemical structures of solid compounds. More importantly for the subject of this chapter, the isotropic chemical shifts are conformation dependent and thus provide information on the geometries of the molecules considered.

3.2.3 NMR Studies of Orientational and Conformational Order in Longitudinal Thermotropic Polymers As argued in Sections 1.7 and 1.8, differences in molecular structures of PLCs cause large differences in their properties. For this reason, we shall survey NMR studies taking one class at a time. We begin with longitudinal polymers, which are the most widely studied and relatively simple to interpret. Our first example is polyesters:

t~,L:p-C>-~-(CH') "-~l, l CH3 CH3 . J

with n= 10 and molecular weights Mn=4000 and Mn=20 000, which have been studied by Martins et al. 1 0 and Volino and BlumsteinY Under the experimental conditions described in those papers and in the 1 T magnetic field of the 1 H NMR spectrometer, these systems exhibit a nematic plus isotropic biphase between the pure isotropic and pure nematic phases. Representative 1 H NMR spectra of the low-molecular-weight polymer as a function of temperature are shown in Fig. 3.1. Whereas no homogeneous alignment is found for the high-molecular-weight sample, the splitting observed in the biphase and nematic phase of the low-molecular-weight polyester reflects the strong alignment of the macromolecules in the magnetic field of the spectrometer. The nematic order parameter was found to vary between 0·69 at the nematic-isotropic transition and 0·84 at the solid-nematic transition. Analysis ofthe 1 H NMR spectra indicates that the flexible spacers align in the magnetic field with a degree of order comparable to that of the mesogenic core, implying that the chain takes on an average a rather extended conformation. A more detailed insight into the behaviour of the spacer has been reached by Samulski et al.,12 who performed 2H NMR in a magnetic field of 5·87 T on a low-molecularweight polyester of the same series having a perdeuterated alkyl spacer. As shown in Fig. 3.2, the aliphatic quadrupolar splittings segregate into two groups which can be assigned to a-methylene and interior methylene units of the spacer. The orientational order of the spacer C2H2 segments is considerably larger than that found in alkyl units of small-molecule liquid


Fig. 3.1. Representative I H NMR lineshapes of the following low-molecularweight polyester:

te-Q-NIN-P-O-C-(CH2l -q II 10 II x o 0 CH 3 CH3

obtained on cooling: (a) isotropic phase at 147"C; (b) nematic + isotropic biphase at 129°C; (c) macroscopically aligned nematic phase at 110°C; (d) transition

spectrum at 85°C; (e, f) solid phase at 78°C and 40°C, respectively.IO.11

crystals. At the same time, both the magnitudes of the observed quadrupolar splittings and the fact that all of these splittings are not identical, unequivocally show that the spacer is not restricted to the all-trans conformation.

Investigations of the polymers having n = 7 and n = 10 and x varying between 5 and 18 have focused on the biaxiality of molecular ordering in these materialsY Different techniques including 1 H, 2H and magnetic measurements have been used. The main result obtained by Esnault et al. is that the increase of the Szz order parameter with chain length is essentially caused by an increase in the biaxiality Sxx - Syy via a decrease of Syy, Sxx remaining constant. In other words, the increase of order corresponds to a progressive reduction of the orientational fluctuations of the aromatic core in the plane parallel to the ester group; the fluctuations in the plane perpendicular to the ester group are not affected. Moreover, an odd-even effect with the number n of methylene units in the spacer, in phase with that of Sm is found for S xx - Syy. The higher order in even polymers may be due to a better molecular packing.


Fig. 3.2. 38·4 MHz 2H NMR spectra of the following low-molecular weight polyester:

to--/Q>-NIN--<C5'>-O-C-(CO~ -~ ~ ",r ~ 10 ~ •

CH 3 CH3

as a function of temperature. Nematic to isotropic transition temperature: 119°C.12

Orientational and conformational order of a linear thermotropic polyester with a perdeuterated alkyl spacer:

-fo C,,D,, 0 -@- TI 0 --@- 0 C,,D,, 0 -@- 0 TI --@-3-o 0

has been investigated by Bruckner et aJ.14,15 The phase transition temperatures of the polymer are:

170'C 210°C Crystal ~ N ~ I

The orientational order parameter of the meso genic core has been derived from the IH NMR dipolar splittings. It ranges from 0·79 to 0·86 as the temperature decreases from the isotropic-nematic transition to 180°C. It indicates a high degree of alignment of the rigid parts of the chain. The orientational and conformational order parameter of the spacer was determined from the values of the 2H NMR quadrupolar couplings. It is consistent with the existence of highly extended chains undergoing conformational jumps that keep the aliphatic sequences close to the average polymer chain axis.

110 Fran,oise Laupretre

2H NMR studies were also made by Furuya and Abe to elucidate orientational characteristics of ether-type longitudinal PLCs with the repeat unit: 16 ,17

and n = 9 or 10. The following phase transition temperatures were determined by DSC:

140'C 200°C 169°C 204°C

Crystal ~ N ~I

(n=9) (n=10)

Deuterium labels have been introduced either in the aromatic rings or in one of the two kinds of flexible spacer. For the n = 9 polymer a transient nematic-isotropic biphasic region was observed around the isotropicnematic transition. A certain fraction of poorly aligned domains persistently remained in the sample with n = 10 over the nematic temperature range, probably because the external magnetic field of 6·3 T was not strong enough to achieve the alignment of all the chains. Results obtained on the alkyl units have been compared to those measured for dimer molecules and interpreted by using a rotational isomeric state analysis. The trans fractions it for each bond of the flexible spacer in the nematic phase of the dimer and polymer molecules are given in Table 3.1, where for comparison values of/t calculated for the isotropic phase at the same temperature are included in parentheses. These values of it correspond to fairly extended conformations, compatible with the preferred orientation of the meso genic cores. The difference between the dimer and the polymer is very small, indicating that the conformation permitted in the nematic phase is similar. The odd--even alternation of the it values along the chain is marked for n = 10. It is weaker for n = 9. The difference between the dimers and the polymers arises in the orientation of the molecular axis in the liquid-crystalline domain. The order parameter is higher in the polymer.

Rigid-chain polyesters with short flexible side chains may present a complex phase behaviour. For example, the polymer with the repeat unit


Table 3.1 Bond conformations, expressed in terms of trans fraction, it, of the flexible spacer of the liquid crystal dimer and polymer molecules having the repeat unit:

[-@-~~~((H2)n~~~CH2) n-~] at a few degrees below the nematic-isotropic transition temperature. Values in parentheses are those calculated for the isotropic state at the same temperature.a

Polymer Dimer

Bond n=9 n=1O n=9 n=1O

C1-C2 0·57(0·36) 0·36(0·36) 0·54(0·36) 0-42(0·36) C2-C3 0·77(0·55) 0·97(0·55) 0·74(0·57) 0·92(0·56) C3-C4 0·51(0·58) 0-48(0·58) 0·50(0·59) 0·50(0·59) C4-C5 0·67(0·58) 0·91(0·58) 0·64(0·59) 0·86(0·58) C5-C6 0-48(0·58) 0·50(0·58)

a After Ref. 17.

and an intrinsic viscosity of 1·28 dl/g has been studied by Falk and Spiess. 18 It has the following phase transition temperatures:

184°C 228°C 2WC Crystal +-----> layered mesophase +-----> N +-----> I

The polyester has been macroscopically aligned in the 7 T magnetic field of the spectrometer by heating to the isotropic state, then slowly cooling through the nematic and the layered mesophase down to the solid state. As shown by a pronounced angular dependence of the 2H NMR spectra, the high macroscopic alignment of the rigid backbones can be frozen in the solid state. 18 Computer simulations of the spectra show that the two carboxylic functions may be either in cis or trans conformation with respect to the local symmetry axis of the terephthalate ring. In both cases, the angle between the local axis of the terephthalate ring and the chain axis has to be very small, indicating the existence of highly extended chain conformations.

The dependence of 13C NMR chemical shifts on conformations has been used by Uryu and Kato to study the solid-state structure of the following thermotropic PLC: 19

II II roo + TC-@-C-o-@-o-CCH216-o-@-For the original sample as obtained by low-temperature solution polymerization, the l3C NMR solid-state spectrum is consistent with an

112 Franroise Lauprerre

all-trans conformation of the hexamethylene spacer in the plane of the phenoxy group. Variations in the lineshape and chemical shifts of both the aromatic carbons artha to the spacer (Fig. 3.3: carbon atoms 2 and 2') and the aliphatic carbons (Fig. 3.3: carbon atoms Ct., fJ and y) show that the conformations of the hexamethylene spacer and ether linkage are a function of thermal history.

Aharoni et al. have combined the above technique with X-ray diffraction to elucidate the conformational nature of the x- and y-alkylene segments in the crystalline and mesomorphic states of poly(esteramides):20

o 0 0 0 i I~II II~II II T\QrC-o-CCH2) y-o-c----B-~-C-CCH2) xc-~

H H

The conformations in the x-alkylene segments have been found to be exclusively trans, whereas the conformations in the y-alkylene segments include both trans and gauche states. The presence of the gauche conformers in the y-segments allows the chains in the crystalline and mesomorphic states to adopt an overall zig-zag geometry. The conformation of the y-segments involving the Ct. CH2-CH2 bond is solely trans for the crystalline polester-amides with y = 2. In the crystalline and mesomorphic states of polymers with y = 3, the conformations about this CH2 bond exist as a distribution of, or rapidly interconverts between, trans and gauche conformers. For y = 4, 5 and 9 only gauche states are observed for this bond.

Some generalizations can be drawn from the above set of studies. It is a common observation that thermotropic longitudinal PLCs with a high molecular weight and therefore a high melt viscosity are difficult to orient and require a high value of the threshold magnetic field strength. Moreover, under certain circumstances probably because of the polydispersity of the sample and the value of the applied magnetic field, a biphasic behaviour is observed. Also, the rate of orientation is much slower in PLCs as compared to their MLC counterparts. A systematic study of the orientation kinetics in a magnetic field was performed for:

/I II II II II II { o 0 ~o 0 0 0 ti C ---{CH2is- c-o-@-o G-{CH215-C-+0-@-~-@-O-+c-@-0 'y

(x~O, y~O, x+ y~ 1, z=x+ y-l)


200 100 ppm

Fig. 3.3. Solid-state 67·8 MHz 13C MAS DD NMR spectra of the following polyester:

1~%7 7 6 ~ -{§},-3 2 -!§F2 3 i c 0 c-o 0 O-C~-::H2-C~-C~-C~-C~-O 0 ls 5 4 I I 4

77 32' 11 ~ ~ 11 23

with different thermal histories: (a) original sample; (b) sample cooled slowly from 280°C; (c) sample cooled rapidly from 280°C; (d) sample quenched in liquid N z

from 280°C; (e) sample quenched in liquid N z from 360°c. 19


by Moore and Stupp.21 For samples varying in molecular weight, the 1 H NMR results exemplify a great sensitivity of orientation rate to the molar mass. For sample with molecular weight 6900, characteristic orientation times range from 47 s at 199°C to more than 104 s at 170°C. At the timescale of the experiments, no significant macroscopic orientation occurs below 170DC, although the most prominent endothermic transition in the DSC trace is around 145°C. Isothermal ageing time of the mesophase before application of the magnetic field is also an interesting variable; it leads to a drastic reduction in orientation time and reflects the non-equilibrium nature of the meso phase structure. The orienting effect of surfaces on PLCs has also been examined.22 Indeed, surface interactions can have strong effects on the alignment of thermotropic polymers when coupled with other force fields.

As also shown in the examples reviewed above, once the orientation is achieved, the orientational order is much higher in longitudinal PLCs than in their MLC analogues. In all the systems studied, the chains adopt highly extended conformations. For polymers having a flexible spacer, the conformations of the different spacer bonds are either all-trans or include a number of gauche states. It is of interest to compare these experimental results with the predictions of conformational energy calculations.23 25 The latter show the influence on the conformational order of the spacer units, of the spacer length (even~odd effect) and the nature of the functional groups which link the polymethylene segments to the mesogenic core. When the spacer is attached to the rigid mesogenic core by an oxygen atom or an O(O=C) group, spacers which have an even number of methylene units exhibit a significant population of highly extended conformations that allow nearly parallel alignment of the rigid cores along the major extension axis. On the other hand, for an odd number of spacer carbons, the number of highly extended chains is much smaller, the mesogenic cores are tilted by an angle of about 30° from the major extension axis and the order parameter is smaller. When the polymethylene spacer is attached to the mesogenic core by a (C=O)O group, the population of extended conformers is reduced significantly for the chains with even-numbered methylene spacers. When compared with those of odd-numbered methylene spacers linked in the same way, they exhibit even smaller fractions of extended conformers. 23

Finally, the fact that highly oriented chain conformations of longitudinal PLCs can be frozen in the solid state is of major interest for the


mechanical properties of melt spun fibres from the nematic phase. In the case of the polyester with the following repeat unit2 6

-tO~CDO--@--OOC ~OCD: -((}J:)s-CD2T

a

and phase transition temperatures

177'C 294'C Tg = 400 C; C rys tal+----------> N +----------> I

analysis of various 2H NMR experiments performed by Muller et al. on as-spun and annealed fibres have shown that they are characterized by a high degree of orientational and conformational order. In addition, practically all director axes are aligned in the draw direction. Annealing results in a better development of crystallinity without further improvement of the degree of order. Discussion of these data in relation to the mechanical properties of the fibres clearly shows a strong correlation between the existence of a high modulus and tensile strength for the fibres and the high degree of orientational and conformational order that is achieved in these materials.

3.2.4 NMR Studies of Orientational and Conformational Order in Side Chain Thermotropic Polymers Unlike longitudinal PLCs, side chain ones orient quite easily in the external magnetic field of the spectrometer.

3.2.4.1 Polyacrylates and polymethacrylates Side chain liquid crystal polyacrylates (PAm,n) and polymethacrylates (PMAm,n) with the general formula:

(Rl = H or CH 3 ) have been extensively studied by Spiess et aJ.2,27 34 by 2H NMR using selective deuteration of either the main chain protons or different side chain ones. The PA2.l and PA6 ,l polyacrylates whose phase transition temperatures are

335K 389K T g +----------> N +----------> I


and

T 308K S 370K 396K I g~ A~N~

have been macroscopically oriented in the 8·4 T magnetic field of the spectrometer through heating into the isotropic phase and then slowly cooling to the liquid-crystalline phases and finally to the glassy state. The anglar-dependent 2H NMR lineshapes of the deuterated terminal phenylene ring have been analysed in terms of an orientation distribution function. This function has been found to be Gaussian with widths of I8'5°C and 10.5 0 for the frozen m=2 nematic phase and the frozen m=6 smectic phase, respectively. Under these conditions, the molecular order of the meso genic cores28-30 is high: 0·88 and 0·65 for m = 6 and m = 2, respectively. With regards to the spacer, the order parameter derived from the behaviour of the methylene group next to the meso genic core is virtually the same as that obtained from the mesogenic core itself. By contrast, the order parameter measured for the methylene group adjacent to the chain backbone of PA6 ,! is reduced to 50% of the value calculated for the mesogenic group,28 whereas the equivalent methylene group in low-molecular-weight analogues exhibits a reduction by only 25%. This indicates a substantially higher fraction of gauche conformers in the spacer of the polymer than in the alkyl chains of low-molecular-weight liquid crystals.

As for the chain backbone, the PMA6 ,4 polymethacrylate has been studied in the frozen smectic phase, The preferential orientation of the C-C2H3 bonds at 54° with respect to the director is only consistent with an orientation of the chain perpendicular to it. The width of the orientational distribution of the C-CH3 bonds with respect to the director has been determined to be only ± 20°, which indicates the high conformational order at this site. 3! However, line shapes obtained from the two de ute rated sites of the main chain cannot be interpreted in terms of an all-trans conformation of the chain backbone. The link between the polymer chain and the mesogens is provided by the conformationally ordered quaternary carbon in the polymer chain. Molecular order of the main chain has also been investigated for PA6 ,! and PMA6 ,!. 32 As for PMA6 ,4, PMA6 ,! polymethacrylate in its frozen nematic phase has a main chain which is oriented perpendicular to the director. By contrast, in the frozen smectic phase of the polyacrylate the chain is found to be oriented parallel to the director. These results are consistent with the molecular shape of the polymer as a rotational ellipsoid with its long axis parallel to the director for the polyacrylate and perpendicular to the


director for the polymethacrylate. They seem related more to the difference in stiffnesses of the polyacrylate and polymethacrylate chains than to the actual nature of the frozen liquid crystalline phase.

High-resolution solid-state 13C NMR experiments were also made for the PA 6 •1 polymer in the frozen mesophase. 2 .34 Information about molecular order has been derived from the chemical shift anisotropy by using two dimensional magic-angle spinning NMR. 35 The principle of the experiment is the following: if one slowly spins an ordered sample with the order axis not parallel to the rotor axis, one observes phase and intensity modulations of the centre bands and side bands. The modulations depend on the position of the rotor at the time the spectrum is excited. Therefore, by synchronizing data acquisition with the rotor position, a side band pattern is obtained for each resonance. The spectrum thus obtained for PA6 •1 cooled below its glass transition is shown in Fig. 3.4. The first-order side band is enlarged, showing the resolution and assignment of the different carbons: main chain, spacer, carbonyl and quaternary carbon of the outer phenyl ring. Subspectral

M=-4--------_~~. __ --_

s

<; ~ 0

«i ~ v ., '0 0 E

'"i _____ molecular slnJcture

Fig. 3.4. 2D 13C MAS DD NMR spectrum of PA6 •1 in the frozen state. The first-order sideband is enlarged showing the assignment of different carbons: (1) main chain carbon atom; (2) spacer carbon atom in f3 position to the oxygen atom; (3) acrylic carbon atom; (4) quaternary carbon in the outer phenylene ring

connected to the OCH3 group.2.34

118 Franroise Lauprhre

analysis was used to determine the orientation distribution functions of all these carbon atoms. The results obtained so far are in general agreement with those derived from 2H NMR for the spacer and the mesogenic core. The highest order is found for the mesogenic group: it is approximately twice the order of the spacer. However, no gradient in the order parameter is observed for the spacer carbons. The gradient from the spacer to the acrylic carbon next to the main chain is weak. There is a strong decrease in order from that position to the main chain itself through only one bond, suggesting the role of the rotation about that bond for the decoupling.

3.2.4.2 Polysiloxanes The fact that the chemical shift anisotropy is a tensor has also been used by Oulyadi et al. to study the orientation of thermotropic polysiloxanes.36 13C NMR spectra obtained in the presence of highpower proton decoupling and in the absence of magic-angle sample spinning have been recorded for comb polysiloxanes (Pn.m) of the formula

<[H3 --\5i-0 ) 35

(IGl2) no-@-o-t----<Q>--o-CmH 2m+! o

and exhibiting a smectic A thermotropism. Proton-decoupled 75 MHz 13C NMR spectra of a static sample of P S,l at different temperatures are shown in Fig. 3.5. They have been obtained after slowly cooling the P S •1

polysiloxane from the isotropic state to the smectic A phase in the strong magnetic field (7 T) of the NMR spectrometer. The spectrum recorded at a temperature of 117°C, just below the SA~I transition temperature, has narrow lines which are strongly shifted with respect to those observed at 129°C in the isotropic phase. Aromatic and carboxyl carbon lines are shifted to lower fields, whereas methylene carbon lines are shifted to higher fields. The lines remain narrow but the chemical shift variations increase with decreasing temperature and then tend to limiting values. The observation of narrow lines with chemical shifts which differ from the isotropic values indicates that the polysiloxane molecules are oriented in the magnetic field. These experiments provide values of the order parameter associated with the mesogenic unit. Variations of Szz as a function of the number of carbons in the spacer, n, and terminal group, m, are plotted in Figs 3.6 and 3.7, respectively. It can be seen from these figures that these two parts of the molecule do not play the same role. The

CH 3

{~i - OJ I 35 CH 2 e I

~H ( d

~H2 c

C;:H2 b

~H2 a o

6~12 5 3

4 o I C-O

6~r&~: ~ OCH 3


129 ·C Isotropic phase

IJJ J. .. I ... UIl,! .... '.1, 1"", 1", 1"T"f4rJ,""'I'",...,.

117 ' C S A phase

260.0 200.0 140.0 80.0 PPM

119

CC,b

d

20.0 - 10.0

Fig. 3.5. 75 MHz 13e DD NMR spectra of a static sample of P 5 1 as a function of decreasing temperature. 36 •

120

Szz

F ranroise Laupretre

1.0,...---------------------

0.8

0.6

0.4

0.2

• D 8

I!! 8

§

o. 0 '------''------'_--''_--I._----L_---'-_----'-_-L_--'-_~ o 10 20 30 40 50

!Tiso - T)

Fig. 3.6. Temperature dependence of the molecular order parameter Szz in the smectic A phase ofthe Pn,1 polysiloxanes.36 0, n=3;~, n=4; D, n= 5; e, n=6; 0,

n=ll.

Szz 1.0 ,..----------------------,

0.8

0.6

0.4

0.2

o t. o

o

o

o

o •

o o o

o o

0.0 L--'--L----'-_.L..-'--L--'"-.L..-'--L--'"-.L..--'--~

o 10 20 30 40 50 60 70

(Tiso - T)

Fig. 3.7. Temperature dependence of the molecular order parameter Szz in the smectic A phase of the P 4,m polysiloxanes. 36 e, m = 1; D, m = 2; ~, m = 4; 0, m = 8.

NMR Studies of Thermotropic Polymers 121

influence of the spacer length on the order parameter is quite weak, whereas increasing the terminal group increases the order parameter to a much larger extent. The same tendency has been found by zH NMR.37 The conformational order of the (CHz)n part has been determined in terms of trans and gauche populations of the different bonds. Results obtained are somewhat different from those observed on polyacrylates and polymethacrylates. The CHz group next to the meso genic core is in a trans conformation; the next one has a high percentage of gauche conformations exchanging rapidly with trans conformations; the third one is again in a trans conformation. This description of conformational order in the spacer is in general agreement with results obtained from conformational energy calculations on model systems.38 With regard to the alkoxy tail, evidence of gauche conformations has been provided by 2H NMR. 37 Moreover, in the liquid-crystalline phase, the existence of fast molecular motions has been noted: among them are the reorientation of the side chain as a whole about the molecular axis of the mesogenic moiety, internal phenyl ring rotations, OCO group motion responsible for the /3z transition of polysiloxanes and conformational exchange of most of the spacer bonds. 36

3.2.4.3 Other side chain mesomorphic polymers The conformational features of poly(y-n-octadecyl-L-glutamate) were

investigated by Yamanobe et al. by high-resolution solid-state 13C NMR with magic-angle spinning and proton dipolar decoupling in the solid and liquid-crystalline state.39 This polymer forms thermotropic liquid crystals (LCs) by the melting of the side-chain crystallites. In Fig. 3.8 we see the respective spectra as a function of temperature. The values of the CO (amide) chemical shift are characteristic for the right-handed a-helix conformation within the temperature range from 27 to lOO°e. The lines of the n-alkyl carbons change considerably with increasing temperature. Peak I corresponding to carbons in trans conformation disappears above 35°C, whereas peak A associated with the presence of rapidly exchanging trans-gauche conformations increases noticeably. Temperature-dependent chemical shifts are also observed for the a-CH 2, /3-CHz and CH3 carbon atoms. All these results corroborate the fact that the n-alkyl side chains take on an all-trans zig-zag conformation in the crystalline state, whereas they are in a mobile state in the mesophase after the melting of the side chain crystallites.

122 Fran~oise Laupretre

co (amide; II COl ester)

27°C /~ i

150 o

Fig. 3.8. 13C MAS DD NMR spectra of poly(y-n-octadecyl-L-glutamate) as a function of temperature. The main chain carbon peaks are extended. 39

3.2.5 NMR Studies of Orientational and Conformational Order in Discotic Thermotropic Polymers As shown by Huser and Spiess,40,41 the liquid-crystalline nature of the main chain disc polymer containing the units

OR

1 OR

OFl I

1IC'"" ~C~O 0 II II 0 O-C1

OR

NMR Studies o/Thermotropic Polymers 123

and the disc-comb polysiloxane with the following repeat unit

OR

OR

OR

OR I 0 I

O-(CH2),,-Si-CH3

OR I

(with R=CH2-CH2-CH2-CH2-CH3 and selective deuteration either of the aromatic or methylene groups) manifests itself in a macroscopic alignment in high magnetic fields. However high viscosities prevent the formation of interpretable textures under the polarizing microscope. In the main chain polymer with selectively de ute rated methylene groups, the analysis of 2H NMR spectra indicates that the aliphatic C1

carbon of the R group has a well-defined trans conformation, which implies that it takes part in the disco tic order. With increasing distance from the aromatic ring, the disorder in R increases as shown by the presence of gauche conformers at the C3 position.

3.3 NMR INVESTIGATION OF LOCAL DYNAMICS IN MESOMORPHIC POLYMERS

One of the main goals of local dynamics studies using NMR techniques is the identification of the motional processes that are responsible for the transitions observed by mechanical measurements in polymers. Therefore, before examining the results of NMR investigations of local dynamics in mesomorphic polymers, we shall consider various transitions of these materials that may be detected by a number of other techniques including DSC and mechanical measurements.

3.3.1 Glass-Liquid Transition and Secondary Transitions Polymers which do not possess sufficiently high chemical and stereochemical regularity do not crystallize. We know that on a decrease of the temperature they undergo a transition from a highly viscous liquid to a solid glass, called the glass transition. Even for polymers with a high chemical regularity, as, for example, high-density polyethylene, the material never achieves a fully crystalline state; amorphous regions coexist


with crystalline regions, yielding semicrystalline materials. These amorphous parts of the bulk polymer undergo a glass transition.

The same situation prevails for PLCs. Depending on their chemical structure, semicrystalline states mayor may not be obtained. However, owing to the occurrence for these polymers of nematic and/or smectic phases, either the whole material or only the non-crystalline parts of the material will yield an ordered glassy state characterized by a glass transition temperature, Tg (ordered). Such a situation is achieved by cooling the polymer slowly from the high-temperature isotropic state. Of course, if no orientation procedure has been applied, a multidomain sample is obtained.

When a very rapid quenching from the isotropic phase occurs, some regions may retain in the rigid solid the disordered structure of the isotropic state, while other regions undergo the mesophase transitions leading to an ordered glassy state. As polymer segments are less constrained in the disordered glass than in the ordered one, samples prepared under such conditions exhibit two glass transition temperatures, the transition occurring at the lower temperature corresponding to the regions where the isotropic structure has been preserved.

Before considering the effect of the occurrence of a glass transition on the assignment of all the transitions observed by spectroscopic techniques, it is worthwhile to recall that the glass transition observed experimentally is controlled by the kinetics of the structural changes of the material; it has a kinetic nature and the temperature at which it is observed depends on the frequency of the investigating technique. Typically, the glass transition temperature is shifted upward by SOC to 10°C when the measurement frequency is increased by a factor of 10. A commonly used relationship to account quantitatively for the glass transition temperature shift is the so-called WLF equation:42

log(f2/fd = - A[Tg(fd- Tg(f2)]/[B + (Tg(fl )- Tg(f2))] (3.5)

where Tg(fd and Tg(f2) are the glass transition temperatures observed at frequencies fl and f2' respectively, and A and B are constants which depend slightly on the polymer but can be assumed as universal, to a first approximation. This relationship can be derived from a free-volume theory. It describes the non-Arrhenian behaviour of the molecular phenomena responsible for the glass transition.

At temperatures below the glass transition (a), most polymers in the solid state exhibit secondary transitions ({3, y, etc.) which correspond to the motions of groups in the side chain or the main chain, such as the {3 transition observed in polyacrylates or polymethacrylates assigned to the


rotational motion of the ester group around its linkage to the main chain. These secondary transitions are characterized by a frequency dependence which corresponds to an Arrhenian activated process. At high temperature and high frequencies, the glass transition and {3 transition may merge.

For PLCs, in addition to the secondary and glass transitions, we have meso phase transitions, which are thermodynamic and occur at temperatures which do not depend on frequency. Thus, for example, lowfrequency measurements on a non-crystalline PLC show with increasing temperature the following transitions: y, {3, glass transition, smectic~ nematic and nematic~isotropic. At higher frequencies, but lower than the frequency, fa, where Tg(fa) is equal to the smectic-nematic transition temperature, TSN , the same behaviour is observed; the glass transition associated with the ordered glassy state can be detected. At frequencies higher than fa, molecular motions in the isotropic, nematic, smectic liquid phases or solid state can be investigated. However, the highfrequency glass transition of the smectic glassy state no longer exists because the smectic~nematic transition changes the original smectic structural arrangement of the polymer segments. These features are very important for assigning the various transitions observed using spectroscopic techniques. In particular, a transition occurring at high frequency in a temperature range where the glass transition is observed by calorimetry experiments does not correspond to the glass transition process but to secondary relaxations.

Finally, it follows from these considerations that in studies of molecular processes involved in the glass transition of the ordered, glassy state using a spectroscopic technique operating at a frequency f, it is essential that the mesophase found in the glassy state remains till a temperature higher than Tg(f). Typically, to perform such a study at 105 Hz it is necessary that the polymer does not undergo any mesophase transition at a temperature lower than Tg(calorimetry) + 50°C (at 108 Hz, it would be Tg(calorimetry) + 80°C).

3.3.2 Principles of the NMR Experiments

3.3.2.1 'H NMR Detailed information on molecular dynamics can be obtained from the determination of the magnetic relaxation times T" T,p and T2 .6 The magnetic relaxation in organic materials usually originates from modulation of dipolar proton~proton couplings, and therefore reflects the dynamics of the internuclear proton~proton vectors. To make an effective

126 Franroise Lauprlitre

contribution to T1 , the molecular motions have to possess a frequency near the Larmor proton frequency, which, depending on the static magnetic field, corresponds to correlation times in the range 2 x 10- 12 to 10 - 5 s. Tz values are affected by the same molecular motions as T1 , but also by the low-frequency motions. The spin-lattice relaxation time in the rotating frame, T1P ' depends on spectral densities at frequencies of the same order of magnitude as the intensity, expressed in frequency units, of the applied r.f. field. Consequently, T1p is sensitive to motions with correlation times in the range 10-4 to 1O- 6s, depending on the strength of the r.f. field.

The second moment of the dipolar-broadened 1 H NMR spectrum can be used as an alternative parameter for Tz. It is in principle invariant with respect to nuclear motion in a solid, provided the statistical distribution of the spins remains unaltered. However, as the rate of motion increases with temperature, the spectrum narrows in the centre and wide skirts develop which become too weak for observation. In practice, therefore, the second moments of the recorded spectra decrease with increase in temperature.

3.3.2.2 zH NMR 2H NMR permits the investigation of molecular dynamics in a frequency domain ranging from 10- 1 to 1010 Hz. Spin-lattice relaxation times yield information in the frequency range of the Larmor frequency, i.e. lOB-

1010 HZ,6 whereas other methods such as quadrupolar ech043 and spin alignment44 sequences allow the characteristics of the motion to be determined in the frequency ranges lOs-lOB and 104-10- 1 Hz, respectively. Like 1 H NMR spectra, 2H NMR lineshapes change in the presence of motion. If the motion is rapid on a timescale defined by the inverse width of the spectrum in the absence of motion, the observed spectrum results from an averaged field gradient. It is worth noting that for motions with frequencies below lOB Hz, zH NMR can discriminate between different types of reorientation mechanisms, such as lib ration, conformational jumps or rotational diffusion.45 For the purpose of testing models of molecular motions, an interesting pulse sequence has been recently developed46 which allows the individual spectral densities for the various magnetic relaxation times to be determined.

3.3.2.3 13C NMR High-resolution solid-state 13C NMR, which allows the observation of one signal per magnetically non-equivalent carbon, makes it possible to


follow the dynamic behaviour of each part of a molecule independently.9 NMR parameters which are sensitive to molecular motions include the different relaxation times as well as the spectrum lineshape, the strength of the dipolar coupling or the chemical shift anisotropy. As for I Hand 2H NMR, the available frequency windows depend on the type of measurement which is performed. They cover a very broad range, from slow processes ("'" 1 0 -I Hz) which, for example, manifest themselves in exchange phenomena, to very fast modes (several hundreds of megahertz) accessible by TI measurements. 6 However, if some of the above spectral parameters are mainly determined by molecular motions, others such as Tip, which probes motion in the range of the intensity expressed in frequency units of the applied r.f. field, are simultaneously defined by spin dynamics as well as by molecular dynamics.47 50 This may introduce some difficulties in the data interpretation.

3.3.3 NMR Studies of Local Dynamics in Longitudinal Liquid Crystal Polymers

3.3.3.1 I H N M R studies For the liquid crystal aromatic copolyester prepared from 4-hydroxybenzoic acid (HBA) and 2-hydroxy-6-naphthoic acid (HNA) with a monomer ratio (HBA/HNA) equal to 74/26, dynamic mechanical relaxation spectra obtained at low frequencies by Yoon and Jaffe and cited in Ref. 51 have pointed out the existence of a low-temperature relaxation at about - 60°C, and of two further relaxations in the temperature range between 0° and 125°C. In order to identify the type of motions involved in each of the observed mechanical relaxations, Cements et al. have determined IH second moments, <~H2>, on an oriented sample of this copolyester as a function of both the temperature and the angle y between the specimen orientation direction and the static magnetic field. 5 I The I H second moments are shown in Fig. 3.9. At low temperature, both isotropic and anisotropic NMR results are consistent with a rigid structure of random copolymer chains arranged on a hexagonal lattice, where the lattice spacing is determined from X-ray data. As the temperature is raised, the NMR signal changes: at 20°C, the isotropic second moment value can be modelled very satisfactorily by assuming a free rotation of the benzene ring residues about the CI-C4 axis. The observed anisotropy at 20°C also broadly resembles that calculated by assuming mobile HBA groups. At more elevated temperatures (140°C), the observed NMR spectra are consistent with rotation of both HNA and HBA moieties.

128

8

~ 6 (!l

/\ N

:r 4 <I V

2-

Franroise Lauprhre

10 30 50 70 90

Y ( degrees;:

Fig. 3.9. 1H NMR second moments of oriented HBAjHNA copolyester at (a) -60°C, (b) 20°C, and (c) 140°c. 51

3.3.3.2 2 H N M R studies Muller et al. have used 2H-NMR to study molecular order and dynamics of PLCs with the general formula: 52

O*D 0 0 II II -~c- 0 (}....c--<Q>-O-CD, -(CH,l, -CD, -(CH,l, -CD, -(CH, l, -CD, -

D D

II ill ill II

II II O*D CI ° -~c- 0 O--C-@-O--CDz-(CHzl7-CDz-

D D

IV v V


The roman numerals refer to five different polymers, deuterated at different sites in the repeating unit as indicated above. The two groups of polymers differ only in the length of the aliphatic spacer, containing ten and nine methylene units, respectively. The transition temperatures of the polyesters I-III are:

The nematic range of compounds IV-V is smaller, extending from 429 K to 543 K.

The T~ and T2E relaxation times have been measured on macroscopically aligned samples over a wide temperature range. The spin-spin relaxation times T2E are most sensitive to motions with correlation times of the same order of magnitude as the reciprocal of the quadrupolar coupling constant, i.e. in the range from 10- 8 to 10- 4 s, whereas the spin-lattice relaxation times T1 depend on motions with correlation times of the order of the reciprocal of the Larmor frequency, i.e. in the range from 10- 8 to 10- 10 s.

Results have been interpreted in terms of several types of motions, including the intramolecular t<--4g± conformational jumps of the C2H2 group, characterized by a correlation time Ti' and the intermolecular motion of the chain as a whole, undergoing continuous anisotropic diffusion, within an orienting potential: TRII characterizes the chain rotation and TR.L the chain fluctuation. As an example, the temperature dependence of these correlation times obtained from polymer II is shown in Fig. 3.10. In the anisotropic melt, the correlation times of the intermolecular motions of the repeat unit (TRII and TR.L) are of the order of 10- 8 sand 10- 7 s, respectively, while the t<--4g± isomerization is even faster. Below the melting point Tm , the sample appears to be heterogeneous as two components are observed in the NMR experiments. The ratio of the rigid to mobile components in the T~ experiments indicates that it is partially crystallized to an extent of about 60 ± 5%. In the crystalline phase corresponding to the rigid spectrum component, the t<--4g ± isomerization is reduced by more than two orders of magnitude and the corresponding activation energy is approximately 6·7 kJ mol- 1, which is comparable to the values observed in crystalline paraffins. On decreasing temperature, the isomerization process in the amorphous phase is gradually slowed down with an activation energy of 15·9kJmol- 1. This higher value could indicate that the t<--4g ± jumps undergone by the C2H2 groups in the amorphous phase are connected with motions of larger units, possibly involving the nearest-neighbouring phenyl ring, whereas in

130 Fran~oise Laupretre

104.-__________________________ ,

-9 10 r----------------,----,---~

2.0 2.5 3.0 3.5 4.0 1 i T x 103 C K -1 J

4.5

Fig. 3.10. Temperature dependences of 'RII(.&)' 'R.L(e) and 'j(V and _) in the polymer II of Ref. 52, the full and open symbols correspond to the mobile and

rigid components observed below the melting point Tm •52

the crystalline phase only local jumps of the adjacent CH 2 groups could be involved. It is worth noting that the activation energy in the amorphous phase is comparable to that observed for -(CH2--CH2--O)units in a polyaromatic ester with a similar structure. 53 Intermolecular chain motions are also observed in the amorphous phase below the melting point. Their slowing down on reaching the glass transition has the typical feature of non-Arrhenian processes encountered in ordinary amorphous polymers, i.e. a temperature dependence of the correlation time which can be described by the WLF equation.

3.3.3.3 13e N M R studies 13C NMR studies53 •54 have been carried out on a polymer with the repeat unit

I ~ t c-


In the solid state, the investigated sample contains a few isolated regions of disordered amorphous isotropic material, dispersed in a semicrystalline continuum. The calorimetric transitions are:

Glass(1) ~ glass(2) ~ C ~ Sc ~ I

It has not been possible to orient the samples in the magnetic field of the NMR machine and all the experiments have been performed on mesomorphic multidomain samples.

The motion of the aliphatic units has been investigated by measuring the strength of the 13C_l H dipolar coupling, <b2 ). For a powder the value of this interaction is given by:

(3.6)

where rCH is the carbon-proton distance. <b2 ) can be deduced from the rises of 13C magnetization in cross-polarization experiments. 55 Indeed, when carbons are strongly coupled to protons, the rises of polarization can no longer be described by an exponential law. In cross-polarization experiments using very short contact times, <b2 ) can be estimated in a very simple way by measuring the contact time t 1/2 necessary to obtain half of the maximum polarization M( Cf)):

J<b 2)=-Jr-In tl/2

(3.7)

For rCH = 1·09 A, tl/2 = 28 f.1S for a rigid CH group and 20 f.1S for a rigid CH2 group. Experimental tl/2 longer than these rigid-lattice values are evidence for a reduction of the l3C_1H dipolar coupling by motional processes whose frequencies are higher than 105 HZ.55

The contact times, t l /2 , necessary to obtain half of the maximum equilibrium polarization in cross-polarization experiments using very short contact times have been determined at 298 K. At the frequencies involved, i.e. 104_105 Hz, the methylene unit adjacent to the terphenyl moiety has a rigid-lattice behaviour, whereas the next-nearest CH2 group undergoes oscillations on the valence cone of approximatey 200 about one equilibrium conformation. The following CH2 group performs oscillations of large amplitude, or more likely jumps between two equilibrium conformations.

The motions of the carbonyl group and terphenyl moiety have been studied by using the properties of the chemical shift anisotropy. Indeed, the chemical shift anisotropy can be partly or totally


averaged by molecular motions. A well-known example is that of an unprotonated aromatic carbon belonging to a para-substituted phenyl ring. For such a carbon, the principal elements of the tensor are parallel to the axes of the frame defined by the CI-C4 axis and the perpendicular to the plane of the phenyl ring. In the rotation about the para axis, the principal element parallel to this axis is not affected by the motion. In contrast, the two components which are perpendicular to this axis are averaged by the motion. The resulting lineshape corresponds to an axially symmetrical tensor and is markedly different from the rigid-lattice pattern. These two situations have been observed for the protonated aromatic carbons of the above polyester. Analysis of non-spinning proton dipolar-decoupled spectra as a function of temperature clearly indicates that below the crystalsmectic C transition, the protonated terphenyl carbon atoms have a rigid-lattice behaviour. Above this transition, a rapid rotation of the phenyl rings about the CI-C4 axis occurs in the meso phasic state. However, such a motion does not alter the mean orientation of the meso genic group.

Molecular motions in liquid-crystalline poly(ester-amides) of the general formula

rOO 0 0 ~II II~II III ~C-o-(CH2) y-o-c-y-~-C-(CH2) xC-~J

H H

have been investigated by 13C TIp determinations in their semicrystalline state at temperatures below the glass transition temperature. 56

For the methylene carbon adjacent to the amide group, Hatfield and Aharoni have shown that 13C TIp is independent of both x and y. By contrast, when the length of the methylene sequence on the amide side is long, there appears an oscillation in the I3C TIP behaviour of the methylene carbon next to the ester site as a function of y. It is difficult to separate the contribution of static spin-spin processes from that of molecular motions in these experiments.47- 5o However, this 'odd--even effect' might reflect the fact that motion of the methylene carbon next to the ester site is more restricted for even values of y. Comparison of NMR results with thermal behaviour would then suggest that liquid crystallinity is coupled to the molecular dynamics of the ester portion and effectively decoupled from motions in the amide and aromatic portions of the polymer chain.


3.3.4 NMR Studies of Local Dynamics in Side Chain Thermotropic Polymers

3.3.4.1 Polyacrylates and polymethacrylates

133

In order to study the molecular motion of the mesogenic core, Spiess et al. have performed 2H NMR experiments on some of the PAm,n mesomorphic polyacrylates28- 3o whose terminal rings have been fully deuterated. Samples have been macroscopically oriented in the 8-4 T magnetic field of the spectrometer through heating into the isotropic phase and then slowly cooling to the liquid-crystalline phases and finally to the glassy state. We will first consider results obtained in the glassy state and then in the mesophase.

In Figs 3.11 and 3.12 we show the observed and calculated 2H NMR lineshapes of the frozen nematic (PA 2,d and smectic (PA 6,d systems, respectively, with the director n at an angle f3 = 37 0 with respect to the static magnetic field direction. As can be seen in these figures, the line shapes and intensities can be fitted remarkably well by considering 1800 jumps of the terminal phenyl ring about the local C2 axis in the two polymers and taking into account a distribution of correlation times of about 2·5 orders of magnitude. This distribution of correlation times most likely reflects differences in molecular packing resulting in a distribution of the activation energy of this local motion.

Results obtained from these NMR studies have been compared to data derived from dielectric spectroscopy of the same polymers.57 In Figs 3.13(a) and (b) are shown the variation of 10g(T;!), where Tc is the characteristic time of the motion, as a function of the reciprocal of absolute temperature, T, for PA2.! and PA6 ,!. The different relaxation processes observed in the glassy state by the dielectric experiments are shown schematically below:

b relaxation f3 relaxation y relaxation

Of particular interest in this comparison is the f3 relaxation, which is ascribed to the reorientation of the side chain about the Carom~C(OO) bond. Taking into account both the f3 relaxation seen by the dielectric technique and the internal rotation of the terminal phenyl ring about its


Fig. 3.11. Observed (left) and calculated (right) 2H NMR spectra of the frozen nematic PA2,l polysiloxane as a function of temperature. no is the mean

jump frequency.30

Fig. 3.12. Observed (left) and calculated (right) 2H NMR spectra of the frozen smectic PA6 ,1 polysiloxane as a function of temperature. no is the mean

jump frequency, 30

100kHz

lOOk Hz

~J~'467 A.906

....--J~,380 ~,093

no

. A~010 ~'042 ~'741 ~O,282

~~'062

Jt015


346 TjK

2 4 8 50 281 237 180

i' n'9 0 , b , I , c

6 • c ••

•• 0 0 • N

I \ • Phenyl-Flip

J: 4

-----.,.

....

01 ~ 2

0 2.22

6

N

~4 '7 ....

01 o 2

50

i

.. 0 (NMR)

1 \ \ I· • • '8 • • , .' 'n .p

I I ., • I

2.89 3.56 4~22 4.'89 1000 K IT

TIK 346 2~1 237 204 , :9 , nl s

, , , , , p , , "'-·0

I • , . , , ' .0

~ , , ·0 Phenyl-Flip •

N • , c (NMR)

I • • I p. 0

,0\ , • • I ,.

n I , :. • I

I I • I • I , I I I I • ,

I , , I I o '2.22 2.89 3.56 4.22 4.89

1000 KIT

5.56

18 o

[§J

5.56

135

Fig. 3.13. Relaxation maps for (a) unaligned PA2 .!, (b) unaligned PA6 .!; ., dielectric measurements; 0, NMR data for the 1800 phenyl flip.57

symmetry axis revealed by the NMR experiments, one is led to the conclusion that the terminal phenyl ring is involved in both the f3 process and the internal ring rotation. 2H NMR of the terminal phenyl ring is expected to be affected by these two modes, whereas dielectric spectroscopy probes only the former through the reorientation of the electric dipole associated with the COO group. These considerations may explain why, although the two sets of data displayed in Figs 3.13(a) and 3.13(b) in the

136 Fran,oise Lauprhre

temperature-frequency domain of the f3 relaxation are relatively close to each other, the motional frequencies derived from the NMR experiments are systematically higher than those measured by the dielectric technique.

In the liquid crystalline phase, the local 1800 jumps of the phenyl ring are augmented by diffusive rotation about the C2 axis by arbitrary angles as evidenced by the substantial narrowing of the 2H NMR spectra. It is interesting to note that these two motional processes can be separated on cooling into the glassy state, where the diffusive rotation is frozen, whereas the 1800 jumps persist as thermally activated local motions. 28-3o

3.3.4.2 Polysiloxanes Molecular dynamics in the non-oriented state of the P3.1 and PS.1 polysiloxanes have also been investigated by high-resolution solid-state 13C NMR.S8 The P 3•1 polysiloxane is amorphous and the PS.1 polysiloxane is semicrystalline. From the lineshape study of the aromatic part of the high-resolution solid-state 13C NMR spectra recorded over a large temperature range, the rotation of the phenyl rings of the meso genic core has been observed and characterized: at very low temperatures, each pair of ortho carbons of the ring yield two distinct lines on the NMR spectrum. At high temperatures, these ortho carbons, which are magnetically inequivalent in the absence of motion, are rendered equivalent by the phenyl ring motions, leading to an unique NMR line. Between these two extreme situations of the slow exchange and of the rapid exchange, the spectrum lineshape is strongly dependent on the rate of the motion in the range 10 - 1_106 Hz.

Selective pulse sequences have been employed to discriminate between the behaviour of the amorphous and crystalline phases. There does not appear any difference in the phenyl flip dynamics either in the amorphous or in the crystalline phase.

The motions of the methylene units of the spacer were studied by determining the contact time necessary to obtain half of the maximum equilibrium polarization in cross-polarization experiments using very short contact times, t 1/2 • As an example, tl/2 values are shown in Fig. 3.14 for the PS.1 polysiloxane. They demonstrate that the glass transition observed by DSC involves the polymer backbone as well as the nearest spacer carbons. With regard to the glassy smectic A phase made by the meso genic cores, its glass transition is reflected by the increase in mobility of the methylene carbon atom adjacent to the mesogenic unit.


120

I. ... 100 .1

I ,. I' I· • ,ju

'n :l.

0 -N 60 ", ... ~

w • • •

10 CJ

• L

... .. ~ • 0 • • • • •

20 • • Wi

O~~--r-~--~~-10 20 40 6 0 r "C ;

80

Fig. 3.14. Temperature dependence of the till contact times of the CH za(.), CHZb.c(O), CHZd(A) and CHle(e) carbon atoms of the P5.1 spacer (see carbon

numbering in Figure 3.5)58

3.4 NMR INVESTIGATION OF SLOW MOTIONS IN MESOMORPHIC POLYMERS

In Section 3.3 we were interested in the local dynamics of mesomorphic polymers in either their oriented or non-oriented state. However, NMR parameters may also be affected by the long-range orientational fluctuations of the director. Two techniques have recently been proposed to study this collective behaviour in polymer meso phases. 59,60

The first procedure was introduced by Martins et al. 59 It consists of aligning the nematic polymer in the external magnetic field of the NMR spectrometer and then following the lineshape evolution of a monodomain which has been suddenly rotated at a small angle to the static


magnetic field at time O. In this way, a number of viscosity coefficients and the ratio k33/kll of the bend to splay Frank elastic constants can be determined. 59 This technique has recently been applied to longitudinal PLCs with the following formula:

t~NLjQ>-~rCH,),-~l , CH3 CH3 I J

with n = 7 and n = 10.61 - 63 respectively. Experiments were made as a function of temperature and molecular weight. They have shown that for the polymer with n = 10 the twist viscosity coefficient, Y I, can be represented to a good approximation in the whole nematic range by

Y 1 = kS 2 Mi3 exp(E/ R T) (3.8)

where k is a constant, S is the nematic order parameter, E = 64±4kJmol- 1 is an apparent activation energy, and f3=6·0±0·3. Whereas the polymers with n = 7 are characterized by essentially the same value of E, they have a quite different molecular weight dependence (f3 = 6·0 for M < 5000 and f3 < 6·0 for M> 5000).

In the method proposed by Van der Putten et al. for studying ultra-slow director rotation in nematic polymers, the NMR spectrum is monitored while the sample is rotated in the magnetic field of the spectrometer. 60 The rotation speed of the sample container is chosen such that the director orientation follows the container orientation, albeit with a certain lag. By synchronizing the data acquisition with the container orientation, the phase lag is monitored for several hours. Frequencies as low as 10- 6 Hz can be measured precisely. By using this technique values of YI have been determined for a side chain polysiloxane. They are in good agreement with those measured by other authors on the same material.

3.5 CONCLUSIONS

NMR spectroscopy, which has long been recognized as a powerful tool for studying small-molecule liquid crystals, is also a most attractive technique for investigating mesomorphic polymers. As clearly demonstrated by the examples reviewed in this chapter, it is possible to get a detailed insight into orientational and conformational order of the


mesophases. In the field of local dynamics, the selectivity of 13C and 2H NMR, which allows the behaviour of the different parts (meso genic core, spacer, side and end groups) to be examined separately, leads to the identification of the observed motional processes. Some relations with the modes that are involved in the secondary transitions have already been established. However, for these local motions, as well as for phenomena which are responsible for the glass transition in PLCs, a complete understanding has not yet been fully reached. In this respect, one must emphasize the interest of combining several techniques such as NMR measurements with different frequency windows, dielectric and mechanical relaxations, and comparing the obtained results. Finally, among the recent developments of the NMR techniques which are of interest for PLCs, a most promising one is the investigation of ultra-slow motions associated with the long-range fluctuations of the director.

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Chern., 1983, 23, 388. 29. Geib, H., Hisgen, B., Pschorn, u., Ringsdorf, H. & Spiess, H.W., J. Am. Chern.

Soc., 1982, 104, 917. 30. Pschorn, U., Spiess, H.W., Hisgen, B. & Ringsdorf, H., Makromol. Chern.,

1986, 187, 2711. 31. Boeffel, C. & Spiess, H.W., Macromolecules, 1988,21, 1626. 32. Boeffe1, c., Spiess, H.W., Hisgen, B. & Ringsdorf, H., Makromol. Chern. Rapid

Commun., 1986, 7, 777. 33. Harbison, G.S., Vogt, V.D. & Spiess, H.W., J. Chern. Phys., 1987,86, 1206. 34. Blumich, B., Boeffel, c., Harbison, G.S., Yang, Y. & Spiess, H.W., Ber.

Bunsenges. Phys. Chern., 1987, 91, 1100. 35. Harbison, G.S. & Spiess, H.W., Chern. Phys. Lett., 1986, 124, 128. 36. Oulyadi, H., Laupretre, F., Monnerie, L., Mauzac, M., Richard, H. &

Gasparoux, H., Macromolecules, 1990,23, 1965. 37. Mauzac, M., Richard, H. & Latie, L., Macromolecules, 1990,23, 753. 38. Samulski, E.T. & Toriumi, H., J. Chern. Phys., 1983,79, 5194. 39. Yamanobe, T., Tsukahara, M., Komoto, T., Watanabe, J., Ando, I., Uematsu,

I., Deguchi, K., Fujito, T. & Imanari, M., Macromolecules, 1988,21,48. 40. Huser, B. & Spiess, H.W., Makromol. Chern., Rapid Commun., 1988,9,337. 41. Huser, B., Dissertation, Universitat Mainz, 1987. 42. Ferry, J.D., Viscoelastic Properties of Polymers, 3rd edn. Wiley, Chichester,

1980. 43. Powles, J.G. & Strange, J.H., Proc. Phys. Soc., 1963,82,6. 44. Spiess, H.W., J. Chern. Phys., 1980, 72, 6755. 45. Spiess, H.W., Colloid. Polymer. Sci., 1983, 261, 193. 46. Beckmann, P.A., Emsley, J.W., Luckhurst, G.R. & Turner, D.L., Mol. Phys.,

1983, SO, 699. 47. Garroway, A.N., Moniz, W.B. & Resing, H.A., Faraday Symp. Chern. Soc.,

1979, 13, 63. 48. Stejskal, E.O., Schaefer, J. & Steger, T.R., Faraday Symp. Chern. Soc., 1979,13,

56. 49. VanderHart, D.L. & Garroway, A.N., J. Chern. Phys., 1979,71,2773. 50. Schaefer, J., Stejskal, E.O., Steger, T.R., Sefcik, M.D. & McKay, R.A.,

Macromolecules, 1980, 13, 1121. 51. Cements, J., Humphreys, J. & Ward, I.M., 1. Polymer Sci., Polymer Phys. Ed.,

1986, 24, 2293.


52. Muller, K., Meier, P. & Kothe, G., Prog. Nucl. Magn. Reson. Spectrosc., 1985, 17,211.

53. Laupretre, F., Noel, C, Jenkins, W.N. & Williams, G., Faraday Discuss. Chem Soc., 1985, 79, 191.

54. Sergot, P., Laupretre, F., Louis C & Virlet, 1., Polymer, 1981,22,1150. 55. Laupretre, F., Monnerie, L. & Viriet, 1., Macromolecules, 1984, 17, 1397. 56. Hatfield, G.R. & Aharoni, S.M., Macromolecules, 1989,22,3807. 57. Vallerien, S.u., Kremer, F. & Boeffel, C, Liq. Cryst. 1989,4, 79. 58. Oulyadi, H., Laupretre, F., Sergot, P., Monnerie, L., Mauzac, M. & Richard,

H., Macromolecules, 1991,24,2800. 59. Martins, A.F., Esnault, P. & Volino, F., Phys. Rev. Lett., 1986,57,1745. 60. Van der Putten, D., Schwenk, N. & Spiess, H.W., Liq. Cryst., 1989,4, 341. 61. Esnault, P., Volino, F., Martins, A.F., Kumar, S. & Blumstein, A., Mol. Cryst.

Liq. Cryst., 1987, 153, 143. 62. Esnault, P., Casquilho, J.P. & Volino, F., Liq. Cryst., 1988,3, 1425. 63. Klein, T., Jun H.x., Esnault, P., Blumstein, A. & Volino, F., Macromolecules,

1989, 22, 3731.

Chapter 4

Dielectric Relaxation in Macromolecular Liquid Crystals

Jozef K. Moscicki Institute of Physics, J agiellonian University, Krakow, Poland *; and Cor

nell University, Baker Laboratory, Ithaca, New York, USA

4.1 INTRODUCTION

The synthesis of polymeric materials incorporating mesogenic (liquidcrystalline) units in the late 1970s, which was pioneered independently by Finkelmann, Ringsdorf and Wendorff in Mainz,l and by Shibaev and Plate in Moscow,2 has been received with great interest by the polymer and liquid crystal communities around the world. Hybridization of two distinctively different classes of substances, i.e. liquid crystals and polymers, produced bulk properties of new materials that were a melange of properties of constituent substances. The applied scientists found them appealing for their multifunctional character, especially of copolymers, which enable 'engineering' of desired properties of a material conferred by different units. While backbone chains provide polymer characteristics like viscoelasticity and good film- or fiber-forming properties, mesogenic units enhance and facilitate imposition of the orientational order of the material by external fields (electric, magnetic, flow gradient or anchoring field at a substrate), conferring at the same time the anisotropy of many physical properties intrinsic to liquid crystals. Further modification of a material can be made by incorporating into a copolymer chain other functional groups with desired properties. The new materials also

* Address to which correspondence should be sent.

143

144 loze! K. Moscicki

attracted the attention of scientists studying fundamental properties of polymers and liquid crystals, since they brought into view completely new problems associated with molecular interactions on the inter- and intramolecular level. It is well known that the most characteristic properties of polymeric materials like fibers, plastics and rubbers depend directly on molecular mobility in these polymers. Similarly, molecular dynamics in liquid crystals direct their applications in contemporary electronic devices. Not surprisingly, therefore, study of correlations between the molecular dynamics and bulk properties is advancing our knowledge and enhancing the technological exploitation of these materials.

Dielectric relaxation (DR) spectroscopy, which is the main topic of the present chapter, is one of the most frequently used spectroscopic techniques applied in the study of molecular dynamics in condensed matter. Dielectric orientational polarization and relaxation are the result of changes in the orientation of molecular dipoles of constituent molecules. In both liquid crystals and polymers the effective molecular dipole moment results from contributions from different groups of bond moments. As a result, changes in the magnitude and direction of the effective dipole moment originate in inter- and intramolecular movements that the molecule undergoes. Because of the different timescales of the motions involved, by changing the time of observation (the frequency of the probing field) one can effectively separate different modes of molecular dynamics. A great advantage of DR spectroscopy over other spectroscopic techniques is that it enables continuous exploration of a large range of timescales of motions, from as slow as 104 s to as fast as 10- 11 s or even faster. In principle, it is possible, therefore, to study using one technique a broad spectrum of motions as a function of temperature, pressure or composition.

4.2 PRINCIPLES OF DiElECTRIC RELAXATION SPECTROSCOPY

When a molecular system is brought into a static external electric field, all charged particles experience forces tending to move them along the field in the appropriate direction. Since these species usually are more or less mobile, they migrate inside the material, leading to electric polarization of the sample. Firstly, there are small displacements of electrons with respect to the nuclei, creating the induced dipole moments comprising the induced electronic polarization. Secondly, in a similar manner, atoms or groups of atoms are displaced relative to each other, leading to the

Dielectric Relaxation in Macromolecular Liquid Crystals 145

atomic polarization. Thirdly, if there are free ions in the sample, they are moved over long distances, and, if they cannot escape from the sample, they cause interfacial polarization. Finally, if the molecules have permanent dipole moments, then the field tends to direct them, resulting in the orientational polarization of the sample. Since dielectric relaxation spectroscopy is mainly applied to dielectrics, which are insulators, in what follows we neglect effects arising from free carriers in the material (ions, free electrons), assuming that the electric polarization is built up only of the orientational and induced atomic and electronic polarization.

Neglecting higher electric multipoles, the resulting electric polarization P of the homogeneous system can be assumed equal to the dipole density:

P=<M)/V (4.1 )

where V is the volume of the sample, and <M) is an average total dipole moment of this volume (from now on the bold letters will denote vector properties, and the brackets (...) denote a statistical-mechanical average). If the material is isotropic, like most of polymers, then <M) has the same direction as the electric field, and the net resulting polarization is proportional to the electric field strength E:

P=xE (4.2)

The proportionality factor X is called the dielectric susceptibility. Equation (4.2) can be rewritten in terms of the dielectric displacement D:

D=E+4nE=(1 +4nX)E=F.E (4.3)

where B is the dielectric permittivity of the material. If a dielectric is anisotropic, like liquid crystals and most solids, the

scalar susceptibility in eqn (4.2) must be replaced by a tensor. Hence, the permittivity must also be a tensor quantity.

For static electric fields the polarization is in equilibrium with the field. However, if the applied field changes in time, then, in general, the polarization is not in equilibrium with it, since the charged particles and permanent dipole moments may not follow the changes.

Let the electric field vary harmonically, i.e. in the complex notation let it be given by E*(t) = EO exp(iwt), where EO is the amplitude and w the angular frequency of variation. For an isotropic system, the time dependence of the electric displacement is also harmonic: D*(t) = D~ exp(iwt - 6w ), although D~ and the phase difference 6w are frequency dependent. Equation (4.3) now becomes

D*(t) = F.*(w)E*(t) (4.4)

146 laze! K. Mascicki

where

(4.5)

is called the complex dielectric permittivity (or the complex dielectric constant); £' is called the frequency-dependent dielectric constant; £" is called the (dielectric) loss factor; and <>", is the loss angle (i.e. tan(<>OJ) = £" Ie').

As already mentioned, each of the contributing parts of the electric polarization corresponds to motions of different microscopic species. In order to observe each of these motions, the experimental time available for the observation must be correlated with the speed of the motion under study. Therefore, as a consequence of different timescales of the motions, the frequency dependence of e*(w) features different dispersion regions, as shown schematically in Fig. 4.1. If the frequency is sufficiently low, all contributions to polarization have enough time to build up and the net polarization is in equilibrium with the electric field. e'(w) is then equal to the static dielectric constant £0, and e"(w)=O.

E'

2

18 109W-

-2

Fig. 4.1. Dielectric constant 8' and loss factor 8" as a function of frequency for a polar material in the condensed state.

The first dispersion region observed on increasing frequency is associated with the orientational polarization. As the frequency is increased from below to above the frequency for possible dipole reorientation, e'(w) decreases substantially and, at the same time, e"(w) exhibits a broad peak. On further increase of the frequency, e*(w) goes consecutively across regions characteristic for the atomic and electronic polarizations. There


are sharp increases followed by a decrease of e'(O)). e"(O)) has narrow peaks which parallel changes in e'(O)).

Since the variation of the complex dielectric permittivity associated with the orientational polarization is the observable property of a material in dielectric relaxation (DR) spectroscopy, we concentrate on the dispersion region due to the orientational polarization. The particular shapes of the e'(O)) and e"(O)) curves in this region are typical for the relaxation process, characteristic for the rotational motions of molecules in the condensed matter, owing to manifold interactions. The complex dielectric permittivity is related to the autocorrelation function of the total dipole moment of the sample, M, by:

(e*(O)) - e", )/(80 -8(0 ) = po[l- iO)Eiw {F(t)}J = Pa Eiw { - p(tn (4.6)

where F(t) = <M(O)M(t)/<M(O)M(O) is the (normalized) autocorrelation function of M and is usually referred to as the macrosco/!k correlation function, and Eiw denotes the Laplace transform (i = J - 1). eoo is the dielectric constant at frequencies well above the dispersion region for the orientational polarization, cf. Fig. 4.1. The proportionality factor Po accounts for the polarizability of the sample and depends on the shape of the sample and on eo and eoc • For an ideal case, when the macroscopic correlation function is a simple decay function

F(t) = exp( -tlto) (4.7)

where the time constant LO is the correlation time, the complex dielectric constant is given by the familiar Debye equation:

(4.8)

When determined from the experimental data, LO is called the (dielectric) relaxation time. Although the Debye equation frequently gives an adequate description of the orientational polarization behavior, it is not applicable for a number of systems: for example, when the orientational polarization arises from different dipole moments in the system, or reorientations of dipoles are anisotropic. These would gIve rIse to a 'spectrum' of relaxation times:

Igk=l (4.9) k

for a discrete distribution of relaxation times, or 00 00

F(t)= fg(r) exp(-t/!) dL, f9(L)dL=1 (4.10)

o o

148 Jaze! K. Mascicki

for a continuous distribution. Equations (4.9) and (4.10) lead, via (4.6), to the following expressions for e*(01):

(e*(01)-eoo)/(eo-eo,J= Lgk/(1 +i01rd (4.11 ) k

and

00

(e*(01)-e oo )/(80-e oo ) = f g(r)/(1 +i01r)dr (4.12)

o

respectively. Unfortunately, it is not always possible to obtain such closed express

ions for 8*(01) and F(t). Consequently, a number of phenomenological descriptions of the complex dielectric constant and macroscopic polarization autocorrelation function have been proposed in the past to aid in the analytical description of experimental data.

For the purpose of data analysis, the complex dielectric relaxation is usually presented in graphical form as an Argand plot of el/(01) vs e'(01). This graphical representation is commonly known as the Cole-Cole plot. For the case of a simple Debye relaxation (single relaxation time, cf. eqn 4.8), the (e', el/) points lie on a semicircle with the center on the e'-axis at e'=(eo-e oo ) and intersections at c'=G oo and e'=Go; see Fig. 4.2. For a distribution of relaxation times, the Cole-Cole plot always deviates downwards from the semicircle on the interval Go - G,,,, i.e. it lies completely within the semicircle. These deviations take different forms, but

("

Fig. 4.2. The Argand diagrams for the Debye equation (D), the Cole-Cole equation (CC), and the Cole-Davidson equation (CD).


very frequently a spread of experimental points on the (e', e")-plane can be described by analytical functions that are empirical modifications of eqn (4.8), i.e. the Cole-Cole,3 Davidson-Cole4 and Havriliak-Negami5

functions:

Cole-Cole: [e*(w) - e",]/(eo -eC(,) = 1/[1 + (iwro)l-.J

Davidson-Cole: [E*(W) - EooJ/(Eo - Eex,) = 1/(1 + iwro)PJ

Havriliak-N egami: [e*(w)-eooJ/(Eo-eoo )= 1/[1 +(iwro)l-.JP

O:(a:(1

0:({J:(1

(4.13)

(4.14)

(4.15)

Comparison between Cole-Cole plots for the Debye, Cole-Cole and Davidson-Cole equations is made in Fig. 4.2. The arc corresponding to the Cole-Cole equation is symmetrical and forms a portion of a circle, the centre of which is below the E'-axis. The corresponding distribution function of relaxation times is symmetric/ although there is no closed expression for F(t) which would give the Cole-Cole equation. The Davidson-Cole arc is a skewed one, and reflects strongly asymmetric distribution of relaxation times. The distribution is peaked at the critical relaxation time ro, with a decaying tail of shorter relaxation times. There is an exact expression for the autocorrelation function leading to the Davidson-Cole equation.

An alternative description of a non-ideal relaxation process commonly applied to polymeric systems is the Fuoss-Kirkwood equation.? It is usually used when only the loss factor can be measured precisely.

Fuoss-Kirkwood: E"( w) = E;;'ax sech( a' In wr 0) 0:( a' :( 1 (4.16)

where E;;'ax is the maximum value of E"(W) and a' is a phenomenological parameter. The corresponding distribution of relaxation times is symmetric, as in the case of the Cole-Cole equation. (a' should not be confused with a in the Cole-Cole equation; eqn (4.16) reduces to the Debye equation for a'=I, while eqn (4.14) does so for a=O.)

Other quite general, but closer to molecular-level, explanations of a non-ideal relaxation were offered by Williams and Watts8 and by Dissado and Hill. 9 •10 Williams and Watts proposed a modification of a single-exponential correlation function (cf. eqn. 4.7):

Kohlrausch-W illiams-Watts: F(t)=exp( -t/rot O:(b:( 1 (4.17)

150 Joze! K. Moscicki

where K is a phenomenological parameter. (The identical relaxation function was proposed for creep by R. Kohlrausch in 1854;11 we believe it should, therefore, be appropriate to refer to the decay function given by eqn (4.17) as the Kohlrausch-Williams-Watts function.) The complex dielectric constant for the Kohlrausch-Williams-Watts function can be obtained with the aid of eqn (4.6). In general, however, it does not lead to a closed expression for s*(w), complicating its use in analyzing the experimental data.

Dissado and Hill assumed, in contrast to the simple Debye process, that the relaxation behavior is a many-body cooperative process. lO As a result they obtained a quite complicated expression for the complex permittivity.

Dissado-Hill: [s*(w)-s",]/(so-s"J=(1 +iW!DHr l

2F 1 [1-n, I-m; 2-n; (1 +iw!DH)-l] x (4.18)

2 F 1 [1- n, I - m; 2 - n; I]

where 2F 1 [ , ; ; ] is the Gaussian hypergeometric functionY The parameters n, m and !DH represent the influence of three different mechanisms in the process. Index n describes the presence of correlations between neighboring molecules and is defined in the range from zero (no correlations) to unity (complete correlation). !DH is the relaxation time of the correlated neighborhood (a cluster). Parameter m reflects the possibility for exchange of molecules between clusters and varies from zero (no exchange) to unity (perfect hydrodynamic motion). Thus, eqn (4.18) for n = 0 and m = 1 corresponds to the Debye process. The value of 0 log[s"/(so -1:",)]/ ologw gives m and n at w!DH~l and at w!DH~I, respectively.lo It is, therefore, helpful to use the 10nsher log S" vs log W plot when analyzing dielectric data with the aid of Dissado-Hill approach; see Fig. 4.3.

In order to obtain information about molecular dynamics from a dielectric relaxation spectrum, the complex dielectric permittivity is related to the correlation function of the electric dipole moment mi

of the ith species and the dipole moment Ml of a small (in comparison with the whole sample) macroscopic volume V surrounding mi. Ml is the sum of permanent dipole moments in this volume M 1 =~: ~ 1 mb

N being the number of dipole moments in the volume. The


UJ

m n

log w

Fig. 4.3. The 10nscher representation of the Dissado-Hills dielectric loss spectrum. eo is the static permittivity; W DH is the principal relaxation rate; nand mare the Dissado-Hill parameters. Arrows indicate variations of the loss curve on

increasing values of the model parameters.

(normalized) microscopic correlation function is then defined by

ri(t)= (mi(O)Mj (t»/(mi(O)M j (0»

= \ mi(O) Jl mk(t») I \ mi(O) ktl mk(O»)

N

(mi(O)mi(t» + I (mi(O)mk(t»

N (4.19)

mf + I (mi(O)mdO» k"i

A difficulty in linking eqn (4.19) with the complex dielectric permittivity lies in the fact that the permanent dipole moment, even in the absence of the external electric field, is under the influence of the local field due to all other molecules, so it experiences an internal electric field different, in general, from the external one, i.e. E'oe =1= E, which cannot be accounted for in an unambiguous manner. In general, e*(w) can be written in terms of f(t) as

(4.20)

where G(w) is the local field factor. The articular form of G(w) depends on the way in which the dipole moment of the sample M (cf. eqn 4.1) and


the dipole moment of the macroscopic volume Ml are related to each other. It is important to note that, because of the local field factor in eqn (4.19), which is frequency dependent, a single macroscopic relaxation time cannot be associated with a single microscopic relaxation time behavior and vice versa. For example, if a single relaxation behavior is assumed on the microscopic level, i.e. ri(t) is a single exponential, then the complex dielectric permittivity calculated with the aid of eqn (5.20) forms a Cole-Cole plot which is inside the Debye semicircle, which clearly suggests some distribution of macroscopic relaxation times. The same goes for the opposite direction, and at best one may say that the single relaxation time To observed experimentally corresponds to a narrow distribution of microscopic relaxation times centered around a value of Tm = To (SO/Sa,)1/(2eo/e x + 1).6 Luckily, G(w) is frequently a slowly varying function of wand in many cases close to unity in value, so the discrepancy between the microscopic and macroscopic relaxation times is very large. For that reason, G(w) is frequently omitted altogether in eqn (4.20).

If the cross-correlations between different dipole moments can be neglected in eqn (4.19), as in the absence of specific short-range interactions between molecules, e.g. in dilute solutions, then eqn (4.20) reduces to

(4.21 )

where C(t)=<m(0)m(t)/m2 = <u(O)u(t), and we have neglected the local field and dropped a subscript on m (u is the unit vector in the direction of m). The autocorrelation function C(t) is commonly called the (normalized) dipole autocorrelation function or, simply, the dipole correlation function. Although the result in eqn (4.21) is obtained for a very specific case, it follows from it that the dipole relaxation (DR)-spectrum carries a significant amount of information about the molecular dynamics of the system. Importantly, the time dependence of C(l) reflects the particular character of the rotational dynamics of a molecular group with which the dipole is associated. To retrieve this information, however, one has to consider carefully different possible situations influencing the rotational dynamics of the dipole: the molecule structure, its symmetry and flexibility; the structure, i.e. the molecular organization and order of the phase under investigation. In doing this one is usually aided by results from other experimental techniques. Only then can one analyze properly information contained in the dipole correlation function estimated from the DR-spectrum.


In analyzing dielectric relaxation spectra in condensed matter, two different concepts of molecular dynamics are of great significance. The first, which is predominantly applicable to liquids, is to consider the liquid as a dense gas with very frequent collisions, the rate of which is so high that the reorientation of a molecule is completely controlled by the rate. Thus, in this case the rotational dynamics becomes similar to the Brownian translational diffusion and it is known as rotational diffusion. The second approach, which is more frequently used to describe dynamics in molecular crystals and intramolecular dynamics,13 emphasizes reorientation of a molecule in the presence of an orientation-dependent potential with well-defined minima. In this case, the molecule resides for finite time intervals in different potential wells, jumping between them from time to time. Since the time of jump is very short in comparison with the time of residence in the well, the concept is known as reorientation by instantaneous jumps. On infinite increase of the number of potential wells and the corresponding decrease of the angular distance between them, this approach reduces to rotational diffusion. 6

In order to calculate the dipole correlation function, let us consider an ensemble of N identical rigid molecules, each possessing a dipole moment m. To describe the orientation of a molecule in space, two coordinate systems are introduced. The laboratory frame of reference (XYZ), which we will call LF, is traditionally defined as having the Z-axis in the direction of the probing electric field. The molecular frame of reference fixed within the molecule (xyz), which we shall refer to as mF, usually has axes chosen along the principal axes of the moment of inertia tensor (or any other molecular tensor). The orientation of the molecule is then given by the orientation of mF with respect to LF, which is determined by a set of Eulerian angles Q == {IX, /1, Y V4 (see Fig. 4.4). The molecular dynamics of a single molecule is then completely described if we know the probability P(Q, t) dQ of finding the orientation of the molecule in the solid angle element dQ around Q at time t.

The stationary state of the system in the absence of an applied electric field is given by the equilibrium distribution of orientations of molecules P(Qo), defined as the probability density with respect to the infinitesimal solid angle dQo. Then, all we need to know is the (conditional) probability of orientation Q of the molecule at time t if it was Qo at t = 0, i.e. P(Q,tIQo,O}=P(Q,tIQo). The probability of finding the molecule in dQ around Q at t is then equal to P(Q, t) dQ= SP(Qo)P(Q, tlQo) dQo dQ.

Whichever is the mechanism of reorientation, the rotational molecular dynamics in condensed matter is a random process and usually can be


z

a. x

Fig. 4.4. Definition of the Euler angles linking the molecular frame (xyz) and the laboratory frame (XYZ). In terms of molecular motions, y describes the rotation of the molecule about its long axis, f3 describes the declination of the long axis from the Z-axis, and 0( describes the precession of the long axis about the Z-axis.

described more rigorously as the stochastic stationary Markov process,15 for which Q forms a three-dimensional stochastic variable. The dynamics of the system is characterized by a time evolution operator 1) which determines the behavior of the non-equilibrium distribution function P(Q, t):

P(Q, t)= - 1)P(Q, t) (4.22)

and the particular form of 1) depends on the mechanism of molecular reorientation and forces (potentials) acting on the molecule. Once P(Qo) and P(Q, t I Qo) are known, one is able to calculate any desired statisticalmechanical average:

< .. .)= f f. .. P(Qo)P(Q,tIQo)dQodQ, (4.23)

in particular, the dipole autocorrelation function <u*(O)u(t). In order to find an explicit form for the orientational probability

distribution function P(Q, tIQo), the function is routinely expanded in


terms of the elements of an orthogonal set of functions in the space of orientation as the Wigner rotation matrices D{m(Q).14 For example, for the symmetric diffusion case:16 18

P(Q, tIQo)= L [(2) + l)j8n 2]a{m(t)D{;(Qo)DUQ) (4.24) jIm

where a{m(t) are the normalized autocorrelation functions of D{m(Q):

a{m(t) = <D{;(Qo)D{m(Q»/<D{;(Qo)D{m(QO» (4.25)

In the case of the asymmetric diffusion, cross-correlations between different Wigner matrices have to be accounted for.16

The probing electric field 'senses' the component of the molecular dipole moment in the direction of this field. One is, therefore, interested in calculating the dipole correlation function of the dipole moment resolved along the Z-axis of LF. It is convenient for this purpose to build the first-rank irreducible spherical tensor from the principal components of u in mF:

and (4.26)

The transformation rules under rotation from mF to LF then give the components of this tensor in LF: 14

1

u( 1,j) = " D 1 ~ (Q)d l,j') mF L.. )) mF (4.27)

j' = -1

so the projection of u on the Z-axis is

(4.28) j' = - 1

If the sample under investigation is in the isotropic phase, then the distribution of orientation in the absence of the probing field is

(4.29)

and with the aid of eqns (4.24), (4.28), (4.29) and (4.23) one can now calculate the autocorrelation function of Uz for the case of symmetric diffusion: 16

<u~(O)uz(t» = f f u~(O)uz(t)P(Qo)P(Q, tlQo) dQo dQ

= {uia~o(t)+ [(u~ + u; )/2][a~1 (t) + ab -1 (t)]}/3

= {ui a~o(t) + u~ [a~1 (t)+ a~ -1 (t)]}/3 (4.30)


where Ul = Uz and Ut = J u~ + u; are longitudinal and transverse to the z-axis components of u. The dipole autocorrelation function <u~(O)uz(t» is, therefore, a linear combination of the normalized correlation functions of the first-order Wigner matrix elements. It is also interesting to note that both principal components of the molecular dipole moment contribute to <uHO)uz(t» separately, so, in principle, it should be possible to study the dynamics of each of them alone.

In the case of small symmetric-top molecules it is convenient to choose the z-axis along u, so eqn (4.30) reduces to

<u~(O)uz(t» = a6o(t)/3

= <D6~(Qo)D6o(Q»/<D6~(Qo)D6o(Qo»/3 =(1/3)exp(-t/rd (4.31)

Thus, in the case of the isotropic diffusion of small symmetric-top molecules the theory predicts a single-exponential decay of the dipole correlation function.

However, if the diffusion is asymmetric and/or the equilibrium distribution function P(Qo) is not uniform, as in the case of liquid crystals, then the mathematical description of the rotational dynamics becomes more complicated, although it still can be developed along the lines presented above. 16,18,19

In order to understand dielectric relaxation in polymer liquid crystals and in liquid crystal polymers it is necessary to review DR phenomena in monomer liquid crystals and in polymers, and this will be done in turn in the next two sections. We begin with liquid crystals.

4.3 DIELECTRIC SPECTROSCOPY OF LIQUID CRYSTALS

In a simplified picture of the molecular dynamics in the isotropic (liquid) phase and in solid phase, it is commonly assumed that in an isotropic liquid the constituent molecules have freedom of translational and rotational mobility. A solid phase on the other hand, is considered as a state with a high degree of molecular translational and rotational immobilization. Thus at the phase transition, on going from the solid phase to the isotropic phase, molecules are 'liberated' and enabled to move. Traditionally, we refer to this as 'melting' of the translational and rotational degrees of freedom (on going from the solid phase to the isotropic phase, e.g. by warming up a sample)


or their 'freezing' (on going in the opposite direction) at the phase transition.

In many substances melting (freezing) of the translational and of the rotational degrees of freedom happens simultaneously. Other substances, however, show the existence of so-called mesophases which are intermediate between the isotropic liquid and the solid state. It is characteristic of the molecular dynamics of mesophases that molecular mobility in some degrees of freedom is severely restricted, while in the others it is more or less free. If molecules are globular like, for example, molecules of the type (CH 3 hCX where X is Cl, Br, N0 2 or CN,20 on increasing temperature from the solid crystalline phase, a phase transition is observed in which at least one of the rotational degrees of freedom is melted although the centers of mass of the molecules remain organized on a lattice. Because of this rotational mobility, the new phase is called the rotator or plastic crystalline phase. On further increase of temperature the sample can show phase transitions between different plastic crystalline phases, finally undergoing a phase transition to the isotropic liquid at sufficiently high temperatures. Materials featuring plastic crystallinity are commonly known as plastic crystals.

On the other hand, substances made up of strongly asymmetric, rod-like or disc-like molecules, see Table 4.1, show the opposite tendency, i.e. the translational degrees of freedom freeze at much lower temperatures than the rotational ones, resulting in the appearance of a new type of meso phase-the liquid crystalline mesophase. The new mesophase, because of the translational mobility of molecules, shows a higher or lower degree of fluidity, but owing to the existing orientational order of molecules demonstrates at the same time anisotropy of bulk properties, which is characteristic of a crystalline state. This combination of properties explains the name 'liquid-crystalline' given to this phase. Frequently, on changing the temperature, a number of distinctive liquid crystalline phases can be observed in succession for the same substance. Substances with liquid-crystalline phases are commonly referred as liquid crystals.

The existence of the liquid-crystalline state of matter is not only a function of temperature. A large number of organic materials show liquid crystallinity in properly chosen solvents. Systems which exist in the liquid-crystalline state in a definite range of temperature are called thermotropics, while the second group is known as lyotropics. Both groups show a rich polymorphism. While mesophases of rod-like thermotropic liquid crystals can essentially be subdivided into two


Table 4.1 Some examples of liquid-crystalline monomer materials

Chemical structure Name

Rod-like

R -O-<~-O- R'

o

Aromatic esters

Azoxy compounds

R-O-O- R' Biphenyl compounds

Rand R' are most frequently alkyl or alkyloxy chains.

Disc-like

Hexa-alkanoates of benzene

R are alkyl or alkyloxy chains.

main classes-the nematic (N) and smectic (S) phases, the disc-like thermotropic liquid crystals exhibit nematic and canonic mesophases. Lyotropic liquid crystals also show a great variety of different mesophases, constituent units of which can be either anisotropic molecules or, subsequently, on change of composition of a solution, associated groups of many molecules. However, if the constituent entities are rod-like macromolecules or if the associated groups of molecules have an elongated form, they often form a nematic mesophase. Since this chapter essentially deals either with polymers incorporating rod-like mesogenic units or with stiff-chain polymers, we restrict ourselves strictly to recalling only the most characteristic


properties of the thermotropic nematic and smectic phases of elongated molecules.

If we approximate the liquid-crystalline molecules by rod-like particles, then the basic thermotropic meso phases can be visualized schematically as in Fig. 4.5. The nematic phase is the least organized of all liquidcrystalline phases and usually appears just below the isotropic phase. In this phase the centers of mass of the elongated molecules have three translational degrees of freedom and thus are distributed randomly as in an ordinary isotropic liquid. However, the long axes of the neighbouring molecules are preferentially aligned with respect to an axis which is called the director, usually denoted by a unit vector n. If the molecules have mirror symmetry, the ordinary nematic phase is observed. It is nearly always characterized by uniaxial symmetry and equivalence of nand - n. For molecules without mirror symmetry, however, the spatial variation of n leads to an helical structure. If one considers a plane perpendicular to the helix axis, then the direction of n is the same in this plane, but it

Fig. 4.5. Schematic representations of the molecular order in some liquidcrystalline phases. Nand Ch are the ordinary and chiral nematic phases, respectively. A and C are the orthogonal and tilted smectic phases. n is the

director, and p is the cholesteric pitch.


continuously rotates as the plane is swept along the helix axis (see Fig. 4.5). A full rotation of n is completed over a distance p called the pitch. However, the helical structure can result not only from a natural tendency of chiral molecules-it can be induced in the ordinary nematic phase by the boundary conditions. In the first case one talks about the cholesteric or chiral nematics, in the latter case about the twisted nematic.

In smectic phases there is, in addition to the orientational order, also a positional order in at least one direction. As a result, the centers of mass of molecules are, on average, arranged in equidistant planes, so the smectic phase is frequently visualized as a layered structure, in general with uniform director orientation within each layer (see Fig. 4.5). A number of different smectic phases have been found and the differences between them consist in (i) the orientation of the preferred direction of the molecules (the director) with respect to the layer normal, and (ii) the organization of the centers of molecules in the plane of the layer.

The orientation of the director can be either perpendicular to the layers or tilted with respect to the layer normal. The most frequently encountered representative phases for both cases are the smectic A and smectic C, respectively (see Fig. 4.5). For molecules with mirror symmetry, an overall uniform tilt of the director is usually observed in tilted smectic phases, owing to the coupling between the directions of n at the neighbouring layers. However, if molecules forming the smectic phase are enantiomorphic, a chiral phase results, i.e. there is some small angle between the tilted director at the adjacent layers, leading to helicoidal precession of the director as one goes from one layer to another. The situation is somewhat similar to that in the cholesteric phase, and we may consider Fig. 4.5 (Ch) as also a visualization of the rotation that the projection of the director on to the layer plane undergoes on moving along the helix axis. One has to remember, however, that since in the smectic phase molecules are restricted to the layers, the tilt orientation changes in a stepwise manner on going from one layer to another.

We have already mentioned that smectic phases may differ in positional order of molecules in the layer. On the one hand, in smectic phases A and C there is no positional order within the layer-we can think of this layer as a 2D fluid; on the other hand, the degree of positional order in the layer in low-temperature smectics, like B, E, G and H, is so high that, with sustained rotational mobility, it would be more appropriate to classify them as plastic crystals than liquid crystals.


Owing to thermal fluctuations, the orientation of the director is not uniform throughout. In order to obtain 'single crystals' or 'monodomains', the orientation of n can be manipulated to a large extent by applying different external fields, e.g. magnetic, electric, flow velocity gradient, or by special treatment of the solid surfaces between which the nematic is contained. As a result, it is possible to obtain samples of the nematic with nearly uniform orientation on n (this orientation is not ideal since the ordering fields are not strong enough to suppress the thermal fluctuations). Consequently, most experiments can be performed on oriented samples, and in what follows we shall assume that the sample considered is a monodomain unless we specify otherwise. As mentioned, the common feature of either nematic or smectic phases is lack of perfect orientational order of the molecules. To describe the distribution of orientations of the molecules about the director, we introduce the laboratory frame (LF) with the Z-axis coinciding with n, and the molecular frame (mF) with z-axis parallel to the long molecular axis. Then the set of Euler angles Q=(et, /3, y) defines the orientation of the molecule in space. Note that according to Fig. 4.4:

- Angle et describes a precession of the long molecular axis about the director.

- Angle /3 describes a declination of the long molecular axis from the director n.

- Angle y describes a rotation of the molecule about its long axis.

Let the (normalized) probability density function j(r, Qo) describe the distribution of orientations of molecules in the sample, i.e. the probability of finding the molecule at position r with the orientation within dQo around Qo(eto,/3o, Yo) is j(r, Qo)dQo. In the isotropic phase the distribution of centers of mass and orientations must be uniform and j(r, Qo) =

p/8n2 , where p is the density. In the nematic phase, the centers of masses are also distributed randomly, thus j(r, Qo)=pP(Qo). Because the nematic (and also smectic A) is uniaxial, P(Qo) is independent of eto and P(Qo)=P(/3o, Yo)/(2n). Furthermore, we can frequently assume that the molecules can be approximated to rod-like cylinders so that the probability density function also becomes independent of ")'0: P(Qo)=P(/30)!(4n2 ).

It can be expanded in terms of the complete set of Legendre polynomials P 1 (cos /30):

(4.32)


Since P,(cos f3) = D~o(fJ), eqn (4.32) can be rewritten in terms of Wigner matrix elements introduced earlier:

(4.33)

Because of the head-tail symmetry of the nematic phase, only even values of 1 appear in eqns (4.32) and (4.33), 1 = 0,2,4, ....

The expansion coefficients S, can be found from the orthogonality relations of the Legendre polynomials:

1

S, = (P,(cos f3o) = f P, (cos f3o)P(o.o) d(cos f3o) (4.34)

- 1

For 1=0 we get So=l and for 1=2, S2=(3COS2 f3o-1)/2. The coefficient S2 = S is known as the microscopic order parameter,

and is used as a measure of the orientational order of the long molecular axes. Should the long molecular axes be aligned perfectly, then S = 1. In the case of the random distribution of the long axes (as in the isotropic phase), S=O.

The existence of the orientational order of the molecules results in anisotropy of most of the physical properties of liquid-crystalline phases. The magnitude of the observed anisotropy of the macroscopic property under investigation depends on the degree of orientational order, i.e. it depends on S. However, the order parameter can only be determined experimentally if this property can be related directly to a tensor property of the molecules. Particularly suitable for determination of the order parameter is the anisotropy of the magnetic susceptibility, since the macroscopic volume susceptibility XV is directly related to the molecular susceptibility X: XV = N (X), N being the number of molecules in the sample. For uniaxial phases the tensor XV has two principal elements xh = XII and xXx = xh = X.l, where XII is measured in the direction parallel to n, and X.l in the direction perpendicular to n. It can be shown that, for molecules with cylindrical symmetry, the magnetic anisotropy, ~X= XII- X.l' is directly related to the order parameter S:

(4.35)

where XI and XI are respectively the longitudinal and transverse components of the molecular susceptibility. (For other methods of measuring S the reader should refer to textbooks on liquid crystals.21 ,22)


Typical values of S observed experimentally vary from about 0'3-0-4 at the phase transition to the isotropic phase, to about 0·8 at much lower temperatures, indicating a significant orientational disorder present in liquid crystals, since the lower values of S correspond to the average inclination of the long molecular axis from n by about 400

, the higher values of S to about 200 •

Liquid-crystalline molecules possess anisotropy of the electric polarizability, and nearly always a significant permanent dipole moment resulting from contributions from different bond moments, see Table 4.1. Therefore, the dielectric permittivity of liquid crystals is also a tensor quantity. Because of the assumed uniaxiality of the system under consideration, there are again only two principal elements of the dielectric permittivity tensor: t:zz=t:1I and t:xx = t:yy = t:.l' Subscripts II and .1 denote respectively the principal geometries of dielectric measurements, i.e. the probing electric field parallel and perpendicular to the director.

It is well known experimentally that dielectric relaxation spectra of liquid-crystalline phases of elongated molecules have rich structure, dependent on how the phase is oriented with respect to a measuring electric field,23 (see Fig. 4.6). For a monodomain sample of a thermotropic liquid crystal the DR spectrum for the probing electric field parallel to n (Elln) is usually composed of two relaxation domains well separated in frequency; one of them is usually in the range of low megahertz frequencies, and the second in the range of low gigahertz frequencies. On an Argand plot, the low-frequency relaxation can be described by the ideal Debye equation, eqn (4.8), and the high-frequency relaxation by the Cole-Cole relationship, eqn (4.13) (cf. Fig. 4.6). For the probing field perpendicular to n (E.ln), the most frequently observed spectrum in thermotropic liquid crystals is a broad single domain at low gigahertz (see Fig. 4.6). Measurements were also performed on un oriented samples and in these cases they showed the presence of two dispersion regions, one at low megahertz, and the other at low gigahertz frequencies. 24

4.3.1 Dielectric Relaxation in the Uniaxial Phase The character of DR spectra in liquid crystals reflects both the symmetry of these phases and the dynamics of molecules under the influence of the intermolecular ordering potential. In order to explain the observed DR spectra in liquid crystals, a number of different theories based on the use of a general expansion of the orientational probability distribution

164 Jozej' K. Moscicki

a

5 9.8 2.3 T=305K

19.3 E*

50 1 0.5

1000 600 10 0~~~5~----~1~O------~15~--~

b E"

0.5

1000 +

E* 1

(a)

E~ T=339 K

1000

O~~~--/----~~----~~--~ 2.5

(b)

Fig. 4.6. Argand plots of dielectric relaxation spectra of the nematic phase of some typical liquid crystals in two principal geometries of measurement. Numbers refer to frequency (in MHz) of the probing electric field. (From Ref. 23).

functions (cf. Section 4.2) have been proposed,t 7 19.25-29 and we shall summarize briefly the main results of these authors' work.

It has been shown by Kozak et at. 17- 19 ,29 that a qualitative explanation of the multidomain character of DR spectra in the uniaxial liquid crystalline phase with head-tail symmetry (nematic and smectic A) can be made without prior specification of the character of the motions involved in the process, i.e. solely as a consequence of the symmetry of


Let us assume that the rod-like liquid-crystalline molecule possesses a permanent dipole moment U(UJ, ud. In order to estimate DR spectra in the two principal geometries, i.e. for the probing electric field either parallel or perpendicular to the director, one has to know the correlation functions of projections of u on the Z- and X -axes of the laboratory frame (cf. Section 4.2):

1

ull =Uz= I D6;,(Q)U~/) (4.36) j' = -I

and

U.L =Ux= I [D~\j'(Q)-DG(Q)]u~/) (4.37) j'= -I

(since the phase is uniaxial, we have freedom in choosing the direction of the probing field (K.Ln) along the X -axis). Since the equilibrium distribution of orientations of molecules in the phase, P(Qo), is known (cf. eqn (4.33)), one can explicitly calculate the dipole correlation functions (uW(O)ulI(t) and (u!(O)U.L(t) with the aid of eqns (4.22}-{4.24):17-19,29

and

where

(uU (O)ulI (t) =1[(1 + 2S)Aoo(t)u~ + (1- S)Aodt)un (4.38)

Aoo(t) = a6o(t)

AOI (t) = [a61 (t)+ aA -1 (t)]/2

A 1O(t) = [a1o(t) + a_1O(t)]/2

All (t) = all (t)+al-l (t) +a-ll (t) +a-l-l (t)

(4.40)

and alm(t) are the normalized autocorrelation functions of D lm(Q), (cf. eqn (4.24)). Equations (4.38) and (4.39) predict contributions to decay of correlation functions (ulI(O)ulI(t) and (U.L(O)U.L(t) from reorientations of both principal components of the molecular dipole moment, although each of them does it in different manner.

Firstly, we note that respective contributions depend on the degree of orientational order. For example, in the case of perfect orientational order (S = 1) we should not be able to 'see' any contribution to (ulI(O)ulI(t) from Ut , and to (U.L(O)U.L(t) from U\. On the other hand, in


the case of the isotropic phase, eqn (4.38) reduces to eqn (4.29), as should be expected.

Secondly, a decay of the contribution to <ulI(O)ulI(t) from UI is given by Aoo(t), which is the autocorrelation function of D60(0.) (cf. eqns (4.40) and (4.25)). Since D 60(0.) depends only on the {3 angle (cf. Table 4.2), UI

contributes to <ulI(O)ulI(t) solely via reorientation of the molecule about the short molecular axis (see Fig. 4.7). Moreover, because of the symmetry of the phase, it must be identified with the end-over-end reorientation of the long axis.

Table 4.2 Wigner rotation matrices D~dlX,p,y)=e-im·d~dP)e-iky. Explicit forms

mk:

-1 o

-1

(1 +cos Pl/2 sinPlj2 (I-cos Pl/2

of d~dP) functions

o

-sinPlj2 cosp sin Plj2

(I-cos Pl/2 -sinPlj2 (1 +cos Pl/2

The decay of the contribution from Ut , i.e. AOl (t), is the sum of the correlation functions of D6l (0.) and D6-l (0.) and, thus, Ut contributes through a combined motion of the molecule about (i) the long axis (angle y) and (ii) the short axis (angle {3)---see Fig. 4.7.

Similarly, we can identify motions involved in the decay of <U,l(O)U,l(t). Since Alo(t) is built up from correlation functions of Dio(o.) and D ~ 10(0.), it describes precession (of angle IX) of the long axis about n coupled with reorientation of the molecule about the short axis (cf. Table 4.2 and Fig. 4.7). Finally, A 11 (t) involves a combination of motions in IX, {3 and y.

With the aid of eqn (4.20) we calculate from eqns (4.39) and (4.40) the DR spectra in both principal directions of the experiment:

and

[st(W)-SII( 00 )]/[sIlO-slIa,] = Gil [(1 + 2S)uf FII(w)+(l-S)u~ Fl1(w)] (4.41)

[S!(W)-S,l(oo)]/[S,lO-S,loo] =G,l[(l-S)uf Fi(w)+(1 +s/2)uf Fi(w)] (4.42)


z

y

x

y

x x

y y

x x

Fig. 4.7. Schematic representation of molecular motions involved in the four principal modes of dielectric relaxation in the uniaxial liquid crystal phase.

where the proportionality factors Gil and G.L account for the sample polarizability, and incorporate the local field effects.

Thus, taking into account solely the symmetry of the phase, up to four different relaxation domains, two in each direction, are predicted by the theory. However, since we do not know anything about the time behavior of the autocorrelation functions of D tm(n), it is impossible to predict from the theory the frequency range in which contributions from a particular F~(w) should show up. To get this information, specification of a particular model of the molecular dynamics in the phase is necessary.

Theories of dielectric relaxation in the nematic phase which considered particular models of the molecular dynamics have assumed the existence


of the (uniaxial) nematic pseudopotential acting on the long molecular axis. 25 27 As a result, it was possible to predict the sequence of the appearance of dispersion regions F}(w) on the frequency scale and identify them with different modes of molecular motion.

Taking as a reference the dispersion region in the isotropic phase (S = 0), for the electric field parallel to n, the Fli (w) region, which is identified with the end-over-end reorientation of the long molecular axis, should appear shifted to significantly lower frequencies, because the long axis passes over the nematic potential barrier. In fact, this is the strongest manifestation of the presence of the nematic potential in DR spectra of liquid crystals. The frequency range for Fl1(w) is predicted at frequencies slightly above the relaxation process in the isotropic phase. It is interpreted as caused by the fast rotation of the molecule about its long axis, which itself fluctuates about the director on a similar timescale. Since the nematic order directs only the long axes of molecules, it does not influence the rotational dynamics about the long axis, which should be the same as in the isotropic phase. However, because the long axis fluctuates inside a limited solid angle about n (the nematic potential well), the apparent rotational diffusion of the axis is faster than in the isotropic phase. 30•31 Both processes lead to the overall dispersion being shifted towards higher frequencies.

In the direction perpendicular to n, because of the uniaxial symmetry of the phase, the DR spectrum is insensitive to end-over-end reorientations of molecules. Therefore, Fi (w) and Fi (w) involve fast rotation of the molecule about the long axis and small-angle motions on the long axis, resulting in a single broad dispersion region in the comparable frequency range as Fl1(w).

Predictions of the theory were confronted with DR spectra of liquid crystals with a single dipole moment. The most spectacular success of the theory was the explanation of the DR spectra of p,p'-alkylcyanobiphenyls (nCB).32.23 The molecules of nCB have a single dipole moment associated with the cyano end group which is practically parallel to the long molecular axis, i.e. Ul = lui, Ut = 0 (cf. Table 4.1). Therefore, in the nematic phase in the direction parallel to n the only low-frequency Debye-like dispersion region observed (cf. Fig. 4.6a), is associated with end-over-end reorientations of the molecules (Fl1(w)). A single high-frequency dispersion region which shows up in the direction perpendicular to n yields a distribution of relaxation times (see Fig. 4.6a). The frequency range of this dispersion is higher than that in the isotropic phase and, thus, it is assigned to fluctuations of the long axis inside the nematic potential well [Fi (w)].


Unfortunately, there are no liquid crystals with a dipole moment which is perpendicular to the long axis, since they would allow us to verify predictions of the theory in the other limiting case, i.e. Ut = lui and Ul = O. At best, as for example in p,p' -diheptylazoxybenzene (see Table 4.1 and Fig. 4.6b), molecules possess a single dipole moment making a large angle with the long molcular axis, i.e. Ul < Ut . This is sufficient to magnify effects from u~ and suppress contributions from uf. As a result, two dispersion regions are observed in the direction parallel to n. The low-frequency Debye-like dispersion still has a significant increment, while the highfrequency one has a smaller increment and a distribution of relaxation times (Cole-Cole arc) (cf. Fig. 4.6b). The former is identified as caused by Ul taking part in the end-over-end reorientations of molecules [Fl1(w)]; the latter is interpreted as caused by U t rapidly rotating about the long axis, and is seen in this geometry of experiment as a result of fluctuation of the direction of the long axis in respect to n [Fl1(w)]. The only dispersion with a slight distribution of relaxation times and large increment observed in the direction perpendicular to n is interpreted as predominantly caused by Fi (w) processes, since (1 + S/2)u~ ~ (1 - S)uf (cf. eqn (4.42)).

Such an unambiguous interpretation of DR in liquid crystals is not always possible. Most liquid-crystalline molecules possess more than one polar group contributing to the total u, i.e. the kernel polar group and/or the side groups (cf. Table 4.1), the side groups usually being involved in internal rotations. Fortunately, these rotations are performed about bonds that are, to a reasonable approximation, parallel to the long molecular axis. The longitudinal component of u is, thus, independent of internal rotations and the theory holds for this component. Transverse contributions to u from each polar group rotate about the same (long) axis. The motion of these contributions can then be accounted for by introducing auxiliary yi angles describing the orientation of u: in the xy-plane of mF (superscript i distinguishes between different polar groups). In general, one has to expect interactions between u:'s, and the cross-correlations between different dipole moments (u;) in the expression for the microscopic correlation function f(t) (cf. eqn (4.18)), cannot be neglected. There is, however, some experimental evidence that the end group rotation is much faster than the rotation of the kernel about the long axis,23.33 and even that the end groups preserve their independence from the kernel mobility down to a solid phase. 34.35 In such a case, the free rotation (motions are uncorrelated) of each transverse component can be assumed, and their contributions would simply sum up in Fi(w)


and FI1(w): Fl(w) = 1:F~(w), where k =.1 or II. However, if there is any evidence of correlations between different u~'s, an appropriate model has to be constructed.

In addition to the qualitative explanation of the appearance and shape of different dispersion regions in nCB, temperature dependencies of dielectric relaxation times observed experimentally were compared with the estimated temperature dependencies of mean dipole correlation times for F1\(w) and Fi(w). For 7CB, very good qualitative agreement was observed32 (see Fig. 4.8(a)). However, when the temperature range of the liquid crystalline phase is sufficiently broad, as in the case of 5CB, the dielectric relaxation time of the low-frequency process observed parallel

N I

8 1 0 ........... --.0 o-o-Q.....!2 ..... ~

.... 0 ... 0 ......

, '~ If , ,0 , ,0

, ° "

6~i-~~~~~~~'~ 2.8 3.0 3.2 3.4

1000K/T

(a)

.$ 4

o

-4L-~3.0~~3.~5--~4.~0--~4~5--~~ 1000 KIT

(b)

Fig. 4.8. The mean relaxation rate of the dielectric relaxation process as a function of temperature for (a) heptylcyanobiphenyl, and (b) pentylcyanobiphenyl. Open circles and the solid line correspond to experimental data, and

broken lines to theoretical predictions. (From Refs 32 and 36).


to the director shows peculiar temperature dependence leading to a strong curvature of In ill vs llT36 (see Fig. 4.8(b)). This type of behavior of In ill vs liT was also observed in other nematics and smectics,37 and cannot be explained by rotational diffusion theory. It was first pointed out by Zeller36 that the temperature dependence of the form shown in Fig. 4.8(b), is similar to that observed for glass-forming systems. Therefore, the end-over-end reorientation of the molecules should not be considered to be governed by the rotational diffusion process but should rather be associated with a process of large rotational jumps of the long axis. Since the latter is a single-particle process, the relaxation time is well defined, in agreement with the Debye-like dispersion observed experimentally. An appropriate theoretical model based on the free-volume concept was proposed by Diogo and Martins38

and predicts

(4.43)

where To is the temperature at which the volume per molecule is reduced to the close-packing limit, and 8 is a proportionality constant. Diogo and Martins found an excellent verification of eqn (4.43) through the analysis of the temperature dependence of the viscosity coefficients.38 A very good fit of eqn (4.43) to dielectric relaxation data was obtained by Benguigi for a number of thermotropic liquid crystals.37 Additionally, eqn (4.43) was "Iso successfully applied to explain the temperature dependence of the rotational diffusion constant for the reorientation of the long axis of the rod-like spin probe in model membranes in the liquid-crystalline state.39 Despite its success, it is unclear how to accommodate such a single-particle model with the substantial information on strong local intermolecular correlations present in liquidcrystalline phases. It is interesting to note a tendency of liquid crystals to show some properties characteristic of glass-forming materials, e.g. polymers.

4.3.2 Ferroelectric Modes in Chiral Smectic C* Phase One of the most spectacular consequences of the chirality of the smectic chiral C* phase, is the appearance of ferroelectricity. If the Z-axis of the laboratory frame is parallel to the helix axis, then the helical precession of the director can be described by variation of its projection on to the layer plane, ~=~xx+~yy, on moving along the helix axis (~is usually referred to as the order parameter40). For small eo tilt angles, the helix can be then described by ~ = eo cos(cp) x + 80 sin(cp) y = eo cos(qZ) x + 80 sin(qZ)y,


where q is the wave vector of the helix (q = 2n/p, p being the helix pitch). The smectic C* has a two-fold rotation symmetry axis orthogonal to n and in the plane of the layer, and the permanent electric polarization, P, can appear in the direction parallel to this axis, provided the molecules have the transverse component of the dipole moment, P = - Po sin(qZ)x + Po cos(qZ)y, where Po is the magnitude of the spontaneous polarization of the system. Thus, the polarization vector spirals together with n along the helix axis, so the net polarization of the bulk sample becomes zero. Therefore, in unoriented samples of the smectic C* one should observe typical dielectric relaxation processes characteristic of the liquid crystalline phase; see above. However, if the chiral phase is oriented, the polarization fluctuations can be detected by the probing electric field perpendicular to the helix axis. This will happen because the probing field disturbs the helix by forcing local polarization to align along the field. As a result, two spectacular low-frequency relaxation processes are observed in the smectic C* forming materials.40 44 The lower-frequency Goldstone mode appears in the smectic C* because of fluctuations in the azimuthal orientation of the director. The higherfrequency soft mode appears in the vicinity of the smectic C*-smectic A transition as a result of fluctuations in the inclination of n from the layer normal.

The appearance of both modes is explained on the basis of the C*-A phase transition theory.40.45 The theory predicts the appearance of a contribution to the complex dielectric permittivity associated with the fluctuations in the tilt angle and polarization, in the vicinity of the C*-A transition.45 Four normal modes describing these fluctuations are predicted to exist in the smectic C*; (two in-phase modes, the inphase orientation (the Goldstone mode) and amplitude (the soft mode) fluctuations of the order parameters; and two analogous outof-phase modes. On transition to the smectic A, since the helix is gone, the in-phase and out-of-phase modes degenerate into one in-phase (the soft mode) and one out-of-phase mode, respectively. The dielectric spectrum is then the sum of different contributions from the normal modes of the system. Numerical calculations have shown, however, that the strength of DR processes resulting from the out-of-phase modes is significantly smaller than from the in-phase modes,45 which might explain why they have not so far been observed experimentally.

Experimentally the Goldstone and soft modes exhibit properties that are unusual for dielectric process.40-43 Although both can be always well


described by the Cole-Cole curve with a small broadening parameter, away from the phase transition the strength of the Goldstone mode, which is associated with the spontaneous polarization, is several times greater than the strength of the soft mode. In fact, because the frequency ranges in which both modes appear overlap, the soft mode is essentially buried in the high-frequency side of the Goldstone peak. On approaching the phase transition, as the helix starts to unwind, the amplitude of the Goldstone mode decreases dramatically, but at the same time the opposite happens to the amplitude of the soft mode. Also the dependence of relaxation time on temperature reflects dramatic differences between processes. The Goldstone mode relaxation time remains long, and is essentially constant within the smectic C* (see Fig. 4.9(a)). The relaxation time of the soft mode which is much shorter away from the phase transition, however, diverges critically (r = k/I T - TeAl) as the phase transition is approached,divergence being significantly faster in the C* than in the A phase (cf. Fig. 4.9(a)).

Precise characterization of the soft mode was thus possible because the helix can also be unwound by strong magnetic fields. 40 .42 .43 This enables suppression of the Goldstone mode and uncovers the soft mode (see Fig. 4.9(b)).

N I "-" -

400.----------------------------,0,---,

320

240

160

~a l 80 ~

Goldstone mode CD

" o

cf9 0 o

o

o

• 5000V/cm

" 8000 Vlcm

o 10000V/cm

o 12000V/cm \ ~f O+------.--~·T-----_,------._----~ 320 324 328 332 336 340

T/K

(a)

Fig. 4.9. Ferroelectric Goldstone and soft modes in the chiral smectic phase. (a) Temperature dependence of the relaxation rate Ie; (b) the dielectric loss spectrum

as a function of bias d.c. voltage. (From Ref. 42, with permission.)


200

160

160

" 120 0

0 2 4 6

E(IOOO 'V/cm) E"

80 326.39K

40 Soft mode

0 2 3 4 5 6 7

log (t 1Hz)

(b)

Fig. 4.9.-Contd.

4.4 DIELECTRIC RELAXATION IN POLYMERS

In the previous section the appearance of liquid crystalline phases was discussed in terms of freezing (or melting) of the intermolecular degrees of freedom of elongated molecules. In this respect, the molecules were considered as rigid entities, i.e. accounting for mobility of different molecular groups in the molecule (internal degrees of freedom) was not necessary for the explanation of the appearance of the liquidcrystalline state. However, when it comes to polymer macromolecules the situation is different. The flexibility of polymer chains, the extent of the chain backbone, and mutual entanglement of chains usually make the dynamics in the internal degrees of freedom far more important in defining the properties of the polymer sample. Flexibility of the polymer chain originates in the ability of chemical groups to rotate about single covalent bonds in the chain backbone and, thus, this rotation is the most important rotational motion active in polymers. Summation of rotations of different groups in the backbone produces an effect of motion of the segments of the chain relative to others, leading, in turn, to the motion of the whole macromolecule. The overall translational and rotational motions of the macromolecules are, therefore, integrated with


the conformational changes (intramolecular rotations) and distortions of the backbone chain, which results in all types of molecular motions being strongly coupled together. In a bulk sample, the situation is even more complicated because the mutual entanglement of the chains introduces steric constraints on segmental and overall motion. Steric obstacles caused by the neighborhood force moving chains to perform a sort of reptilian random walk that is a combination of translation and conformational changes, through a network of the other chains. Molecular dynamics in polymers is, then, a complicated phenomenon requiring careful experimental and theoretical studies.

Because of the entanglement of polymer chains, bulk polymers do not show a natural tendency to form a perfect crystalline phase on cooling from the isotropic melt. At best, in the solid phase there exist regions of crystalline and amorphous material, but frequently the solid phase takes the form of the amorphous glassy state. As in liquid crystals, the appearance of different states in a bulk polymer sample and properties of these states can be associated with cessation (or onset) of molecular motion. For example, in the isotropic melt, macromolecules possess freedom of both translational and segmental rotational movements. On cooling the melt of an amorphous polymer, a transition to a rubbery or a leathery state is observed. This transition is associated with a cessation of large-scale translational motions, but the segmental mobility remains active. On further lowering of the temperature, the glass transition is observed, below which the segmental motions become frozen. It has to be noted that the glass transition is not the same type of transition as the crystal melting or rubber melting transitions. In the latter case, transitions are thermodynamic transitions and the cessation of particular modes of motion is instantaneous at the transition, and accordingly the temperatures at which they are observed do not depend on the frequency of the method used for their investigation. The glass transition, on the contrary, is associated with continuous slowdown of the molecular mobility on cooling of the sample until the motion is virtually discontinued. Thus, the transition is kinetic in nature and the temperature at which it is observed depends on the frequency of the measurement technique. The glass transition temperature-frequency relationship can be explained in terms of the free-volume concept, already mentioned in the previous section (cf. eqn (4.43)).

If a polymer material exhibits the crystalline solid phase, then on cooling of the melt a direct transition from the melt to crystalline solid is observed. However, as we mentioned above, the solid phase is not a


perfect crystal phase, but a heterogeneous mixture of crystallites and amorphous material. It is, therefore, difficult to associate this transition with the cessation of all molecular motions. The translational motion will be frozen out but some segmental motion will persist in the amorphous regions and probably also in crystallites. Therefore, on further decrease of temperature one would expect to observe the glass transition in the amorphous material and possible additional transitions in crystallites. In fact, a number of subsidiary transitions are observed in solid polymers and most of them are associated with freezing/melting of different intramolecular modes of motion.

Further complexity is added to the picture of molecular dynamics by the presence of side groups attached to the backbone chain which can move independently from the backbone, or by the possibility of network formation by the cross-linked chains.

A polymer chain usually incorporates many polar groups, contributions from which produce the effective dipole moment of the macromolecule. The direction and the magnitude of this dipole moment varies as a result of intramolecular rotations of polar groups and overall reorientation of the macromolecule. DR spectroscopy therefore becomes a very valuable experimental method of studying the motion of the polymer macromolecules. In order to obtain detailed molecular interpretation of the DR spectra of polymers, the dipole correlation formalism function is commonly used. 46 The microscopic correlation function f;(t) (cf. eqn (4.18)) for the polymer chains should be considered separately for each different kind m; of dipole moment present on the chain, and the intrachain and interchain cross-correlation terms must be accounted for. To estimate the intrachain cross-correlation terms in an unambiguous manner, it is important to know how the unit dipoles are fixed to the chain, and a number of different situations are possible in practice. For example, the dipoles can be attached parallel and unidirectional to the chain backbone, producing an effective dipole moment which is parallel to the end-to-end vector. Information given by the DR spectrum is, in this case, limited to the dynamics of the end-to-end vector. If the chain is flexible then the cooperation of the motion of constituent dipole moments is accounted for in terms of normal mode analysis,47,48 but when the chain is stiff it is more appropriate to discuss the dynamics in terms of rigid-body rotation, i,e. extremely cooperative motion of the unit dipole moments. If the dipoles are parallel but their directions are random, then there is no correlation between the end-to-end vector and the resultant dipole moment, and dielectric results are interpreted in terms of intra-


molecular rotations of the constituent dipole moments, and a study of cooperation (cross-correlations) in the motion of different unit dipoles is possible. A similar situation occurs in the case of the unit dipoles perpendicular or fixed at some angle to the backbone contour. Finally, if the dipoles are fixed on the side chains, rotation independent of the backbone is usually considered.

The interchain cross-correlations would show up first in bulk samples below the melting temperature, when the translational mobility of chains is severely restricted. In the rubbery, glassy or crystalline states the situation becomes somewhat similar to that encountered in mesomorphic materials and a detailed knowledge of the intermolecular organisation of the chains is important for proper construction of ri(t).

Since most of the polymer properties significant to technological applications must depend on molecular motions present in the material, detailed studies of molecular dynamics of the polymer chains are inevitably very important. In order to understand the molecular dynamics of polymer chains in bulk, one customarily begins with estimation of the intrinsic flexibility of the chain. This can be done best when the intramolecular effects are isolated from the interchain forces and effects. For this purpose, studies are performed on polymers dissolved in carefully chosen solvents. Apart from solvation effects, this offers the best available approximation of the 'free' chain. Knowing the intramolecular dynamics of the isolated chain, one can move to investigate polymer mobility in a bulk sample, in particular to clarify intcrchain phenomena.

4.4.1 Polymers in Diluted Solutions When a polymer solution is diluted to the extent that interactions between different polymer chains can be neglected, then the dipole autocorrelation function ri(t) (ef. eqn (4.19)), aside from the autocorrelation term, accounts only for cross-correlations between the given dipole mi and the rest of the n - I dipoles on the same chain:

<mi(O)mj(t) + I <mj(O)mdt) ri(t)= _ _ __--'k"_i ____ _ (4.44)

mi + I <mj(OlmdO) k"i

On the one hand, the more flexible the chain, the more independent are the rotations of the molecular groups in the chain, and there is less and

178 Jozer K. Moscicki

less cross-correlation between different dipole moments and the correlation function tends to fi(t) = <mi(O)mi(t)/mf. On the other hand, the stiffer the chain, the more important the intrachain cross-correlations become. In the limiting case of the rigid chain, fi(t) reduces to the correlation function of the total dipole moment of the chain m = ~mi' i.e. f m(t) = <m(O)m(t)/m2.

4.1.1 Flexible chains One of the important problems that was studied with the help of DR spectroscopy is the influence of substituents on the dynamics and flexibility of otherwise similar chains. For example, if different substituents with permanent dipole moments perpendicular to the backbone are chosen, then the dielectric relaxation time should vary as a function of the size and steric hindrance of the substituent. Such a comparative study for vinyl and related polymers in different solvents clearly showed that the relaxation time decreases as the substituent group becomes larger and bulkier, i.e. the increase in steric interactions in the backbone chain leads to decreased flexibility.49 Such an effect was even more pronounced for acrylic polymers in toluene. In this case, the dipole moment serving as a probe is associated with the ester carbonyl moiety. On the one hand, direct substitution of the hydrogen by the ethyl group on the backbone creates a dramatic decrease in the dielectric relaxation time by nearly two orders of magnitude. 5 0 52 On the other hand, elongation of the flexible chain in the ester group had a slight effect, which should be expected. 50

4.4.1.2 Stiff chains Completely different aspects of molecular dynamics are studied in dilute solutions of polymer chains when a polymer backbone chain is stiff, since it is the whole-molecule rotation that is investigated. The rigidity of the chain results from the presence of relatively high energy barriers opposing segmental rotation. In the hypothetical case of infinite energy barriers the chain should form a helix, leading to an overall rod-like shape of the macromolecule. However, since the barriers are finite, any real 'rod-like' macromolecule must have some degree of flexibility. When the chain is short, deviation of the shape from the rod-like is negligible, but on increasing the chain length even slight curvature can build up to cause an overall change of the shape to coil-like. Since the flexibility of the chain increases with


temperature (the probability of segmental rotation increases), the same effect can be observed in the temperature variation, i.e. a chain which is rod-like at low temperatures can undergo the rod-coil transformation at sufficiently elevated temperatures. Good examples of such stiff chains are poly(n-alkylisocyanate) and poly(y-benzyl L-glutamate) compounds or DNA.

The molecular dynamics of a polymer chain in the rod-like and in the coil-like conformations is dramatically different. Kirkwood and coworkers predicted that the rotational relaxation time of a rod in a dilute solution should vary approximately as the cube of the rod length. 5 3 Since the length of a macromolecule in the rod-like state is proportional to the number of monomer units and, thus, to the molecular weight, the end-to-end vector autocorrelation time should vary as the cube of the molecular weight. Poly (n-alkyl isocyanate) (PAIC) polymers are particularly suitable for such studies. It was established that the spatial configuration of PAle chains consists of helices that can be pictured as close to the cis-trans conformation. 54 The unit dipole moments are rigidly attached to the backbone chain. Because of the helical structure of the polymer chain, the total dipole moment of the macromolecule is parallel to the helix axis and its magnitude is proportional to the helix length (molecular weight). If the helix is rigid, the DR studies should reflect this fact by showing the relaxation time as proportional to the cube of molecular weight, and the dielectric increment as proportional to the square of molecular weight. Therefore, any deviation of the macromolecule from the rigid-rod form should be clearly seen in the DR spectra. Such behavior has indeed been observed for dilute solutions of fractionated samples of poly(n-butylisocyanate) (PBIC).54 For molecular weight below 105, the relaxation time dependence on the molecular weight follows the cube law. However, for molecular weight greater than 105, the dependence of dielectric relaxation time on the molecular weight deviates gradually from the cube law towards a three-halves power law,s4 the latter being characteristic for the coil-like conformation of the chain (see Fig. 4.10). Results for poly(y-benzyl L-glutamate)55 and for DNA 56

confirm these observations.

4.4.2 Flexible Polymers in Bulk On increasing the concentration of a polymer in solution, different chains begin to collide with each other, and as the concentration increases further and the frequency of collisions rises, the chain mobility becomes dramatically restricted. The situation of a chain in a mesh of more or less

180 ]ozej' K. Moscicki

-1r-----.-----,-----~~

;&SLOPE=i.5 / ~

/

-2

pi _ -3 et° :q, SLOPE=27(.,

~ -4 rn o

-5

I -6

-7~ __ ~ ____ ~ ____ ~~ 4 5 6 7

log Mw

Fig. 4.10. Rod-like polymers in dilute solution. The molecular weight (Mw) dependence of the mean relaxation time, To, for poly(n-alkylisocyanate) compounds. (From Ref. 54, with permission of the American Institute of

Physics.)

entangled neighbors can be envisaged as the inside of a hypothetical worm-like tube defined by the topological constraints imposed by the other chains. If chains are flexible, then size of the tube allows the chain a limited change of conformations, and the chain can move translationally in a snake-like manner (reptation).5 7

At the same time, the segment rotational motions also become severely limited. Steric collisions of the chain with the neighbours cause, at first, departure of the segment orientation correlation function from a single exponential decay function characteristic of a free chain, 58 and on further increase of concentration the segment dynamics change completely.

When studying the molecular dynamics of a polymer chain in a concentrated solution or in bulk material, both the intra- and intermolecular interactions have to be considered, and the correlation function r;(t), as well as the autocorrelation term, should account for


cross-correlations between the given dipole mi and N dipoles that it interacts (correlates) with (cf. Section 4.2):

N

<m;(O)mi(t) + L <mi(O)mdt) fi(t)= N b'i (4.45)

m? + L <mi(O)mdO» k#i

i.e. not only for correlations between the dipoles on the same chain but also for those on the other chains in the neighborhood.

The more densely packed the chains, the more limited become the conformational changes, and the more restricted are reorientational motions of different polar groups on chains. Somehow, similar to the situation in mesomorphic materials (cf. Section 4.3), at sufficiently low temperatures some or all of the various component polar groups can be immobilized. On increasing temperature, 'melting' of motion in different rotational degrees of freedom can occur at various temperatures, thus giving rise to the distinct transition and relaxation processes. It has to be stressed that because of the complexity of the dynamical structure of polymers, the term 'melting' is used here in a very general manner. A change in intramolecular mobility may include not only liberation of a group to rotate, but also, for example, change in the magnitude of the orientational freedom, or change in the degree of cooperativity of neighboring groups.

An enormous amount of work has been carried out in the study of dielectric relaxation in solid polymers. It is well documented that most dipolar amorphous polymers exhibit multiple dielectric relaxation behavior. 59 •6o When only one kind of dipole moment is present in the chain, then at low temperatures most often two processes, named a and 13, are observed, the a-process being the lower-frequency one. As the temperature is increased (at constant pressure) the frequency of the a-process increases more rapidly than that of the f3-relaxation, so that at higher temperatures the two processes merge into a single aI3-process, as shown schematically in Fig. 4.11. Note, that while a plot of InUc) vs liT, is linear (Arrhenius-like) for the I3-process and exists below the rubber-glass transition, it is strongly non-linear for the a-relaxation, reflecting cessation of this process at the rubberglass transition.

The general properties of a, 13- and aI3-processes have been extensively studied over a wide range of temperature, frequency and pressure, and

182

N I

CJl .Q

Jaze! K. Mascicki

I

~

1000 KiT

Fig. 4.11. Typical temperature dependence of the relaxation rate for the rx and f3 processes in material forming a glassy state. Tg is the glass transition temperature.

for polymers of different chemical structure and stereoregularity. The most significant findings can be summarized as follows:

The intramolecular motions which relax the total dielectric increment, L\s=so-s'"" or in other words the mean square dipole moment (m2 ),

can generally be divided into two classes: (i) fast local motions and (ii) large-scale, slower Brownian motions of chains. The former are responsible for the [J-process, and the latter for the a-process. The mean square dipole moment, (m2 ) = m2 + L(mi(O)mk(O), where the sum is taken over all mk that correlate with the given dipole mi (cf. eqn. (4.45) and Section 4.2), is then partially relaxed by local motions, and the rest via the Brownian motions.

The character of the local motions ([J-process) depends on how the dipoles are attached to the polymer chain. They can either be rigidly connected to the backbone chain (e.g. polysiloxanes) or be bonded to the flexible side chains (e.g. polyalkylmethacrylates). In the former case the [J-relaxation must involve the backbone chain motions; in the latter case the mean square dipole moment may be partially relaxed via the


reorientation of the flexible side groups. In both cases, the local molecular motions are possible only if the local environment allows for it, and this provides the common factor for the processes in the two types of polymer. This strong dependence of the local dynamics on the environment makes the f3-process a highly cooperative one, leading to a broad dielectric spectrum.

The high-temperature af3-process has the same mechanism as the a-process. As the temperature is raised, the a-process rapidly becomes faster than the f3-process, becoming then the sale process, which relaxes all available <m2 ). Consequently, the magnitude of the af3-process corresponds to all of <m2 ). It should be noted, however, that the particular mechanism of the af3-process is not the same for all polymers, but depends on the structure of the polymer chain. For example, in lower polyalkylmethacrylates at high temperatures, the af3-relaxation is governed predominantly by the Brownian motions of the backbone chain. On elongation of alkyl chains, however, i.e. in the higher polyalkylmethacrylates, the Brownian motions of the constituent dipoles are significantly controlled by the dynamics of alkyl side chains.

The a- and af3-processes are characterized by a broad asymmetric dielectric relaxation spectrum, which can be well represented by the Kohlrausch Williams-Watts (KWW) decay function (cf. eqn. (4.17)). The major factor leading to the broad DR spectra for a- and af3-relaxations is that chain segments relax in cooperation with their environment. In order to explain the mechanism of this relaxation, the concepts of defect diffusion and free-volume fluctuation are used. For example, Bendler61 has proposed a model in which the KWW function is interpreted as the survival probability of a frozen segment in a swarm of hopping defects with a stable waiting-time distribution At - K for defect motion.

A qualitative theoretical interpretation of the experimental findings on the basis of the dipole correlation function has been proposed by Williams and Watts. 62 They assumed that a reference dipole mi may find itself in a number of different environments and, consequently, the fraction of <m2 ) relaxed by local motions will depend on the particular environment. The environments may, in turn, be changed by micro-Brownian motions. Let the probability of finding the dipole in an environment I be p(l). If the fraction of the mean square dipole moment relaxed by the f3-process is bmf = <m2 )d<m2 ), then (l-bmf) of the mean square dipole moment is left unrelaxed by this

184 laze! K. Moscicki

process. Assuming that the dipole decay function for the fi-process is rp/(t), and for the:x-process is r.(t), one can write the effective dipole correlation function as

ri(t) = r ,(t)· I [(I - bmf) + bmf r P/(t)] p(l) (4.46)

where summation extends over all possible environments. Equation (4.46) offers a qualitative explanation of the main features of dielectric relaxation in bulk polymers. Firstly, at low temperatures r P/(t) decays faster than r,(t), so two separated relaxation domains will be observed. The high-frequency domain associated with a range of rp/(t)'s, each contributing p(l) 8mf to the effective magnitude of the process, will feature a broad dielectric spectrum. The ratio of the relative magnitudes of the :x- and fi-processes should be Ip(l)[(1 - bmt )/[Ip(l) bmfJ, i.e. an increase in the magnitude of the fi-process will be accompanied by a decrease in the magnitude of the :x-process, and vice versa. Secondly, at high temperatures the environments relax so rapidly that the local motions are not effective, i.e. ri(t) ~ r,(t), (cf. eqn (4.46)), and the :xfi-process has the mechanism of the :x-process.

Finally, it should be noted that the fi-relaxation is not the only local motion process observed in amorphous polymers athigh frequencies and at low temperatures. If different kinds of dipole moments are present in a chain, then other fast processes may show up reflecting the local motions of different kinds of dipoles, and customarily they are labelled with the consecutive greek letters; ,-, b-, ... and so on.

4.4.3 Rod-like Polymers in Concentrated Solutions Changes in the dynamics of rod-like polymers in solution on increasing concentration are in some respects similar to and in others dissimilar to those observed in flexible chains. Initially, as concentration rises, the rod-like chains bcgin to collide with each other but, nevertheless, still can perform end-over-end reorientation, the rate of which is mainly defined by the solvent. On further increase in concentration, the collisions become so frequent that a chain finds itself temporarily trapped in a cage formed by its neighbors. It can still freely diffuse translation ally along the long axis, but translational motion perpendicular to the long axis and reorientation of the axis are restricted to the space defined by the cage. In particular, the solid angle LlO available for the in-solvent reorientation of the long axis is smaller than that available in dilute solution, LlO < 4n (the whole solid angle). However, since constraints relax frequently, the chain


moves from one trap to another, and on a longer timescale the whole 4n-space is accessible to the long axis.

Additional increase in concentration reduces the size of the trap and, thus of ,10. Now the chain finds itself residing most of the time in the cage, which may change orientation only from time to time and only by a small angle, as a result of fluctuations in the neighborhood. The time during which the long axis can complete end-over-end reorientation increases enormously. In fact, at sufficiently high concentration the size of the trap may reach the dimensions of the chain, which then becomes completely immobilized. This may be manifested macroscopically by precipitation of the polymer from the solution or by gellation of the solution. For some rod-like polymers in properly chosen solvents, before the critical concentration for precipitation or gellation is reached, a transition from isotropic solution to anisotropic solution is observed,63 in which the chains show the orientational order characteristic of the nematic phase. The rotational dynamics of the chains in nematic solution will be considered in the next section.

To stress the importance of the rigidity and shape anisotropy of rod-like macromolecules for their dynamics, Doi6466 proposed a model in which the real system is simplified to a system of hard, thin monodisperse rods. The distance p between a rod and the nearest neighbor is taken as a good measure of the size of a cage that the rod is trapped in. The solid angle ,10 available to the rod for rotational diffusion is then of the order of (p/L)2, where L is the rod length. A change of the trap takes place if the rod diffuses translationally along its axis a sufficient distance to escape, or if the restricting neighbor does the same, creating a 'defect' in the cage. These two processes define an average time of residence in the trap: I1tcxL2/D\ where D' is the translational diffusion constant for diffusion parallel to the long axis. Since the small change of equilibrium orientation gained via the escape-from-trap mechanism is only of the order of 110, M is also the time between 'jumps' from one cage to the next. The rotational relaxation time for large-scale reorientation is then estimated to be of the order of M/l1ft

The theory predicts two relaxation processes for monodisperse rods: a fast one associated with the wobbling of a rod in the cage, and a slow one associated with the overall reorientation of the rod. The fraction of the mean square dipole moments relaxed in the former process should be of the order of 2110 and the relaxation time should be much shorter than that characteristic of the unrestricted diffusion

186 J ozeJ K. M oscicki

in diluted solution. For the latter, the relaxed fraction is expected to be (1- 2Ml), and the rotational dependence of relaxation time on the rod length should change on going from a dilute solution to a semidilute solution from the cube law to a seventh power law-i.e. very dramatically. At the same time the relaxation time should vary as the square of rod concentration.

It should be noted that the two-step rotational dynamics of the long axis in a semidilute solution qualitatively resembles the f3- and aprocesses in the solid state of flexible polymers; the short-timescale motion in the cage may be compared to the f3-process, the long-timescale fluctuations in the environment with the a-process.

The fast and slow processes have been observed experimentally using Kerr effect measurements of PBLG solutions.67 However, since 2~n is a very small quantity, the diffusion in the cage is undetectable by DR spectroscopy and only the slow process can be studied by this method.

Unfortunately, DR studies of the molecular dynamics of rod-like chains in semidilute solutions are very sparse and were performed only for some of the PAles in toluene. 6B 70 As mentioned earlier, PAles are particularly suitable for verification of theoretical predictions by means of DR spectroscopy. The total dipole moment of a PAle chain is parallel to the long axis of a helix formed by the chain. Moreover, the magnitude of this dipole moment is proportional to the length of the chain. Since the dielectric increment and dielectric loss factor maximum are proportional to the mean square dipole moment and the dipole concentration, measurements of each of the quantities provide additional information on the rigidity of the chains on the one hand and on the extend of the solid angle available for the long axis on the other. Verification of the theory via DR spectroscopy would then be straightforward if only PAle samples were not polydisperse. The polydispersity effectively camouflages the dynamics of each individual rod-like chain, and numerical simulations are needed in order to compare the theory with DR experiments. Despite this complexity, dielectric studies have shown that the rotational dynamics in semidilute solutions is dramatically retarded in comparison with a dilute solution (see Fig. 4.16). It was also found that the effective relaxation time varied as the square of polymer concentration in semidilute isotropic solution, in agreement with the theory. The dielectric increment and the maximum of the dielectric loss were linear in concentration, which supports the assumed rigidity of the PAle chain in the range of concentrations studied.


4.5. DIELECTRIC SPECTROSCOPY OF LlaUID CRYSTAL POLYMERS

187

Liquid crystallinity of polymers was first observed more than fifty years ago for a solution of rigid rod-like polymers of biological origin, namely for the rod-like molecules of tobacco mosaic virus dissolved in a properly chosen solvent. 71 ,72 Later a similar phenomenon of lyotropic liquid crystallinity was also discovered for a number of synthetic stiff-chain polymers (see Refs 63 and 73 for reviews), including poly amino acids, wholly aromatic polyamides, polyisocyanides, polyalkylisocyanates, and other homologues, As was demonstrated theoretically by Onsager 74 and Flory,75 the observed lyotropic nematic liquid crystallinity is a consequence of the chain stiffness and of large overall shape anisotropy. When the axial ratio of the rod-like chain exceeds some critical value, steric repulsions between the rods alone are sufficient for the formation of the anisotropic (nematic) order in highly concentrated solutions.

Later it was discovered that semi-rigid polymers in which a series of rigid elements are connected by flexible points can also demonstrate lyotropic liquid crystallinity if the rigid units are sufficiently long. For example, Aharoni 76 studied lyotropic liquid crystallinity in rigid-flexible copolyamides in H 2S04 or DMAc/5%LiCI as a function of the length of the rigid repeat unit (by varying the number of p-benzamide residues), and found that the shortest axial ratio of the rigid element for which liquid crystallinity is observed is of the order of 3-3.5. Although this value is smaller than the minimum axial ratio of 6.417 predicted by Flory for the formation of the nematic phase by rod-like molecules,7 5 it is comparable to the axial ratio for small molecules of the thermotropic nematogen PAA.

Most of the lyotropic liquid-crystalline polymers do not show thermotropic behavior but decompose or undergo other transitions without melting. As an exception, thermotropic liquid crystallinity was observed for some PAICs at sufficiently high temperatures, although at these temperatures they also showed a tendency to decompose rapidly.77 The first semi-rigid polymer (polyester) to exhibit thermotropic liquid crystallinity was reported by Roviello and Sirigu in 1975,78 and, subsequently a number of thermotropic polymers have been synthesized. 79 84 As in the case of lyotropic semi-rigid chains, a common feature of these chains is the presence of rigid elongated units separated by flexible spacers, e.g. several -CH2- groups. It turned out from initial studies that the critical minimum elongation of the rigid segment in


thermotropic polymers should be of the order of 4. Since the axial ratios of most of the monomer liquid crystalline molecules are equal to or larger than this value, it was natural to expect that polymer chains incorporating mesogenic units would demonstrate liquid crystallinity.

The first liquid crystal polymers (LCPs), in which structural moieties known to lead to mesomorphic behavior in small-molecule liquid crystals have been incorporated into the main chain, were reported in early 1980s.85-88 The presence of the mesogenic units in the polymer chain enhances the tendency for the material to form the liquid crystalline state, and these polymers also show better thermal stability. Both the nematic and smectic phases can be observed in these materials,87 (see Fig. 4.12).

Because of the structural similarity between semi-rigid thermotropic polymers and LCPs, i.e. the presence of rigid moieties separated by flexible spacers, all polymers of this kind are commonly referred to as the main chain liquid crystal polymers (mcLCP).

The mesogenic units can also be attached as pendants to a flexible backbone through flexible spacers (see Fig. 4.13), resulting in a completely new class of polymers with unique properties, so-called the side chain liquid crystal polymers (SCLpC).8991 scLCP manifest a unique competition between the tendency of the mesogenic units towards anisotropic liquid-crystalline order and the tendency of the backbone chain towards isotropic chain conformation. The liquid-crystalline order influences the polymer chain conformations and vice versa. The flexible

Nematic Smectic

n

Fig. 4.12. Schematic representation of molecular order in the (a) nematic and (b) smectic phases of a main chain liquid crystal polymer. n is the director.


(a)

(b)

(c)

Fig. 4.13. Schematic representation of order of the mesogen units and of the backbone chain in (a) the isotropic, (b) the nematic and (c) the smectic phases of a side chain liquid crystal polymer. Om and Db are the directors for the mesogen and

backbone chains, respectively.

polymer backbone chain has a tendency to adjust from a three-dimensinal isotropic conformation in the isotropic melt to an anisotropic conformation in a liquid-crystalline phase. As a result, a number of different new kinds of liquid-crystalline phases may occur for scLCPs depending on the type of liquid-crystalline order (nematic, smectic) present, interactions of the backbone with side chains, and the stiffness of the main chain and of the spacer.

For example, the interplay between the backbone chain and the meso genic side chains can make the nematic phase of scLCPs more complex than a normal nematic (cf. Section 4.3). The nature and magnitude of such an interlay may be expected to depend, to a large extent, on the length, stiffness and shape anisotropy of the otherwise weakly interacting spacer group linking the side chains to the main chain. In the


nematic phase the orientational ordering of the side chain mesogens would influence the orientation of the backbone segments via the spacer group and direct interactions between the backbone and the side chain, thus inducing on average orientational ordering of the main chain. The nematic-like ordering of the backbone could be either parallel or perpendicular to the mesogen side chains,92,93 depending upon the details of the interactions (see Fig, 4.13(b)). This can be rationalized, in general terms, by introducing two director fields for the mesogen group and backbone chain, Om and Db, respectively, with Sm and Sb associated with them as order parameters (d, Section 4,3). Sm behaves then as in ordinary liquid crystals, i.e. it is positive when the side chains are oriented parallel to Om

and negative when the mesogens are perpendicular to it. Sb is positive when the backbone chain has a prolate shape and negative when it is oblate. (Note that purely for symmetry reasons, depending on relative orientation of Om and Db and the signs of Sm and Sb, up to eight different kinds of nematic phase (four uniaxial and four biaxial) may be realized by scLCPs).

In the smectic phase of scLCPs, interactions between the liquidcrystalline and the backbone chain moieties takes a much more dramatic form. The side chains form the strongly interacting part of the smectic layers, while the weakly interacting backbone chain completes the layer and is basically excluded from the strongly interacting part. In a welldeveloped smectic phase the main chain thus finds itself very nearly confined to a two-dimensional plane within the smectic layers, with occasional interlayer hopping (see Fig. 4.13(c)).

Essentially, any desired architecture of the liquid-crystalline polymer can be envisaged and obtained (cf. Chapter 1 by Brostow), although up to now most DR studies have been carried out only for the two most commonly encountered types of the LCP, i.e. the main chain LCP and the side chain LCP.

The motivation behind synthesis of polymers incorporating mesogenic (or rigid) units was different in the cases of the main chain LCPs and the side chain LCPs. This becomes evident if we recall that main chain liquid crystallinity was sought in order to obtain exceptionally strong and stiff fibers without a need for dissolution in very special and corrosive solvents for processing (vide Kevlar and other polyaromatic chains of its class). On the other hand, side chain liquid crystallinity was desired for the expected ability of such polymers to be rapidly oriented by external fields while preserving this orientation on removal of the driving field, this producing a quasi-stable optically anisotropic medium. Clearly, mcLCPs are princi-


pally intended for use in the solid state, and scLCPs in the liquid crystalline state. The character and strategy of the molecular dynamics studies performed to date reflect this situation. As a result, it is impossible to review and discuss these results, and dielectric relaxation in particular, in a compact manner. Therefore, the DR results for each class will be reviewed separately.

4.5.1 Lyotropic Polymers A particular feature of the lyotropic liquid crystallinity of rod-like polymers is that on the concentration-temperature phase diagram the low-concentration pure isotropic phase (cf. Section 4.4) and the highconcentration nematic phase are separated by a significant transition region in which different phases coexist94- 96 (see Fig. 4.14). We mentioned earlier that it has been shown that lyotropic behavior is due predominantly to rigid-rod properties of macromolecules. That is, a solution of rod-like particles of sufficiently high axial ratio (length-to-diameter ratio) should separate spontaneously into two phases, one isotropic and another anisotropic-nematic, as a consequence solely of the rigidity and shape anisotropy of the solute particles (athermallimit). 74,75 To explain the character of the phase diagram in Fig. 4.14, in addition to accounting

80

~ 40 L L+L.C

I-

o

-40

0·2 0·3

v~

Fig. 4.14. Typical concentration-temperature phase diagram for rod-like polymers in concentrated solution. Poly(n-hexyl isocyanate) (PHIC) in toluene. Land

L.c. indicate isotropic and liquid-crystalline states of solution, respectively.99


for repulsions on contact it is necessary to assume the existence a nematic potential, which each rod-like macromolecule experiences.97 In Section 4.4.3 we discussed the influence of the needle-like character of macromolecules on their mobility in concentrated isotropic solutions. The reorienting rod-like particle takes part in two motions; local wobbling inside a small-solid-angle cage defined by neighbors, and hopping from one cage to the next. It is the latter process which ultimately leads to the reorientation of the long molecular axis into the whole of 4n-space, and it is realized on a timescale very much longer than the timescale characteristic of the reorientation in diluted solution.

When the nematic phase is formed, the nematic order of the rod-like molecules imposes a constraint on the reorientation of the long axis of a chain. Because of the extreme elongation of the needle-like macromolecule, the end-over-end reorientation becomes virtually impossible even on a very long timescale, and for most of the time long axis of the chain resides inside some solid angle around the nematic director. The eventual DR-process associated with the end-over-end reorientation should be then pushed down to such low frequencies that it will be practically unobservable in the frequency domain. (Extremely long relaxation times for the overall reorientation were observed in the Kerr effect studies of PAle in toluene. 98 However, the rotational diffusion of the long axis inside a solid angle around the director is a relatively fast process and was clearly observed in DR studies of PAle solutions in toluene.74 76

Extensive DR studies of rod-like polymers in solution covering the isotropic, biphasic and nematic states of solution have been carried out only for two different PAles in toluene: a 1/1 copolymer of n-butyl- and n-nonyl isocyanate (PBNIC)75 and homopoly(n-hexyl isocyanate) (PHIC). 74, 76 The relaxation process was studied as a function of both concentration and temperature. Only one relaxation process, with a broad distribution of relaxation times, was observed in the isotropic and nematic phases up at relatively low frequencies (10- 2 -105 Hz) (see Fig, 4,15). The most important results are summarized in Fig, 4,16, All three characteristic dielectric parameters-the dielectric increment ~F., the maximum of the loss factor £;;', and the logarithm of the mean relaxation rate fe-undergo significant changes across the isotropic-biphasicnematic concentration range,

The most striking result is that all parameters in the pure nematic phase are a few times smaller in magnitude than in the highly concentrated isotropic phase. In the isotropic phase both ~e and e;;' are


30

~:/~ 20 40 60 80 100 E'

(a) (b) Fig. 4.15. Typical Argand plots of the complex dielectric permittivity of rod-like poly(n-alkyl isocyanate) solutions in toluene at room temperature. (a) PHIC and (b) copoly(n-butyl, nonyl-isocyanate) (PBNIC). Land L.c. indicate isotropic and

liquid-crystalline states of solution, respectively, and! is the frequency.70

VO p

Fig. 4.16. Concentration (v~) dependence of fe, fle, and 1>;:' at room temperature for (a) PHIC and (b) PBNIC in toluene. 7o

proportional to the average square dipole moment of a macromolecule, <m2 ), and the number of dipole moments in the sample (= number of macromolecules) (cf. Section 4.4.3). Since the transition to the nematic phase is achieved after concentration is augmented significantly, from that


fact alone one would expect an increase of LlB and B;;'. The dramatic decrease in the magnitude of both quantities on going to the nematic phase suggests, therefore, that only a portion of (m2 ) must have relaxed to this phase. Similarly, one would expect the relaxation time to increase with concentration, but instead acceleration of the relaxation process is observed.

To explain these results, Moscicki proposed a qualitative model of the dielectric relaxation in lyotropic LCPS.99 A main feature of the model was the assumption that each rod can only undergo spatially restricted rotations (hopping between cages) around the nematic director. To rationalize this situation, the theory of small-step rotational diffusion in a conel00.l0l was applied, i.e. it was assumed that the free space available for rotational diffusion of a rod is restricted to some cone coaxial with n and characterized by the cone angle /30, which is a function of the degree of orientational order of rods in solution. For the monodisperse system of rods, the dipole correlation function for such diffusion is

r m(t) = m2 { £5? + £5~ exp[ - v~(v~ + 1 )DtJ + £51 exp[ - v1(vl + l)DtJ} (4.47)

where £5?=(1 +cos /30)2/4, £5~=(l-cos /30)2/12 and £5l =(2-cos /30-cos 2 /30 )/3, and the dependence of v: on /30 can be calculated numerically. 101 D is the diffusion coefficient characteristic of unrestricted, isotropic reorientation of m. It follows from Eqn. (5.47) that a fraction m2 c5? of the total (m 2 ) that remains unrelaxed by a motion in the cone depends strongly on the cone angle. In the limit of unrestricted free diffusion (i.e., /30 = n), c5?=O and the whole (m2) is relaxed. Alternatively, as /30 decreases, less and less (m2) would be relaxed. For sufficiently small /30 the second term in eqn (4.47) is negligible and the relaxation process becomes described by only £5l and vi. For /30 decreasing below n/3, v1 begins to increase rapidly (cf. figure 2. of Ref. 31]. Thus, as the space available for random motion decreases, both the relaxation strength, m2c5L and relaxation time, rl = [vl(1 + vDDtr 1, decrease markedly, as observed experimentally. It is interesting to note that, despite apparent significant difference in the dynamics of a rod-like polymer in solution and of a monomer thermotropic nematogen, the motion of a rod in the cone can be identified as the limit of the F 1 process (cf. Section 4.3, and figure 4 of Ref. 99).

The fact that a real system is more or less polydisperse complicates the application of the theory to experimental data. Detailed numerical calculations have shown, however, that the theory correctly explains all the main features of the DR process in PAIC solutions.99


4.5.2 Thermotropic Polymers Ordinarily, on decreasing temperature, polymers undergo a transItIOn from a highly viscous isotropic fluid to either a non-crystalline solid glass or a semicrystalline solid, depending on chemical and stereochemical regularity (cf. Section 4.4). However, if a polymer exhibits liquid-crystalline behavior, then the isotropic state and a solid state are separated by nematic and/or smectic phases. Consequently, one can encounter a transition from the liquid-crystalline phase to a liquid-crystalline solid glass or a semicrystalline liquid-crystalline solid in these materials. Moreover, if the isotropic phase is very rapidly quenched, a solid state may result with a multidomain structure where some regions retain the isotropic disorder while the other regions demonstrate the order of the liquid-crystalline state.

The appearance of liquid-crystalline order in neat liquid-crystalline polymers should have an influence on the chain dynamics. However, since molecular architectures of mcLCPs and scLCPs significantly differ from each other, the extent of this influence is different for each type of PLC.

Let us first consider the dynamics of the mcLCP. In this case the meso genic rigid units are constituent elements of the main polymer chain. They take part in the segmental backbone dynamics, which do not differ significantly from that of ordinary polymer chains. The only difference may arise from the appearance of liquid-crystalline segmental order in the liquid-crystalline melt and glass. In the absence of any theoretical work on a change in the dynamics of the rigid meso genic units on going from the isotropic to the liquid crystalline melts, the following considerations will be more intuitive than rigorous. For simplicity of discussion we limit our attention to the homopolymer case. In a liquid-crystalline phase of mcLPC, the nematic or smectic structures are formed by a local arrangement of rigid units in the neighboring chains, (see Fig. 12). Firstly, an important observation on the dynamics of these units is that, in comparison with monomer liquid crystals, the end-over-end reorientation of the individual mesogen group, the reorientation so characteristic of monomer liquid crystals, is impossible. In fact, very similar dynamical restriction is also present in the isotropic phase. In both phases, the long axis of the rigid unit would probably at best perform a sort of a space-limited diffusion inside some solid angle defined by a constraint from the chain and the neighboring chains. Approximating this solid angle by a cone, the autocorrelation function of the unit vector along the long axis is identical with that given by eqn 45-47. It follows from the discussion in Section


4.5.1, that the difference in the dynamics of the long axis in the cone on going from the isotropic to nematic (smectic) phase may be notable only if the liquid crystalline orientational order significantly changes the size of the cone. This would lead either to an acceleration of the relaxation process and a decrease of the relaxation strength when the cone is reduced, or to a slowdown of the relaxation rate and an increase of the relaxation strength on expansion of the cone.

Such a change can, however, be small and difficult to detect. In such a case, verification of the existence of this effect can be provided by investigating rapidly quenched samples, where the isotropic and anisotropic regions should coexist (see above). Since the cone size, in general, should depend differently on the temperature for the isotropic and liquidcrystalline melts, significantly different glass transitions for the isotropic and anisotropic regions may be observed. For example, the frequencytemperature dependence of the glass transition in the liquid-crystalline regions should be order parameter dependent,39 while that in the isotropic regions is not (cf. Section 4.3).

The dynamics of the rigid segment about the long axis should not be expected to change noticeably with the transition from the isotropic to the liquid-crystalline melt. Flexible spacers in the polymer chain ensure freedom for the rigid segments to rotate about their long axes without significant constraints. Since this motion is nematic potential insensitive, one should not expect to observe a notable change of the reorientation rate on going from the isotropic to anisotropic state. Therefore, from the point of view of this mode of motion, one should expect similar dynamic behavior to that of ordinary semi flexible chains.

The scLCP dynamics are substantially different, especially those of the mesogen units. The meso genic units are attached to the backbone chain as side groups via a flexible spacer. This leaves them sufficient degree of freedom to perform all reorientations characteristic of monomer liquid crystals. In the isotropic phase, the orientation of the mesogen units is randomized and the rotational dynamics of these units should resemble those of monomer liquid crystals in the isotropic phase. In the nematic or smectic phase, the units are ordered according to the character of the liquid-crystalline phase. Rotational dynamics should, therefore, posses all characteristic features of the monomer liquid-crystalline phase (cf. Section 4.3). The only significant differences in the dynamics in the isotropic and liquid-crystalline phases, in comparison to the monomer liquid crystals, should arise from the fact that the mesogenic units are anchored to the main chain, since this couples their dynamics to the motion of the


backbone chain, and leads to a different mechanism for reorientation about the short axis, in particular end-over-end reorientation (cf. (Fig, 4.13).

Two possible dynamical effects that might be characteristic of LCPs, should be mentioned. Firstly, because of the side chain-backbone coupling, some changes in the dynamics of the side chain may be expected at the smectic-nematic phase transition, in particular when the main chain undergoes conformational transition from 2D configuration in the smectic phase to 3D configuration in the nematic phase. Secondly, in some systems the glassy transition of the backbone chains and of the mesogenic units may not coincide, and two distinct frequency-temperature curves could be observed.

Summarizing the above, we note that main chain polymer liquid crystal dynamics retain far more of the character of the ordinary polymer dynamics than of liquid crystal dynamics. In contrast, the dynamics of side chain polymer liquid crystals demonstrates a number of features characteristic for monomer liquid crystals. The character of the DR studies of LCPs done so far reflects this observation. Because their relatively low shear viscosity in melt makes them significantly easier to process, mcLCPs are natural competitors to glass-fiber-reinforced thermoplastics as self-reinforcing plastics in potential applications. Therefore, the main interest of DR studies has been concentrated on the solid state, mostly in order to identify the relaxation processes in the glassy state observed via mechanical relaxation. scLCPs, in turn, have potential technological importance in optoelectronic devices, non-linear optics and information storage, in which the liquid-crystalline nature plays the most important role (see Chapter 5). Thus, DR spectroscopy has predominantly been used to investigate the liquid-crystalline phases of these materials. When discussing DR results for mcLCPs and scLCPs, we will then follow these approaches.

4.5.2.1 Main chain liquid crystal polymers In the mcLPC chain all permanent dipoles are located on the backbone chain and as such they take part in the inter- and intramolecular motions. In the light of what was mentioned above, we should not expect a DR picture significantly different from that observed for ordinary semi flexible chains. In particular, observation of the influence of orientational order on the dynamics would be problematic, since the end-over-end reorientation of a mesogenic unit is excluded, and other modes of motion should have comparable relaxation frequencies (cf. Section 4.3).


Dielectric measurements of mcLCPs are very sparse and were essentially performed as auxiliary to mechanical studies, reflecting principally interest in the mechanical properties of new materials rather than in peculiarities of the molecular dynamics caused by the liquid crystallinity.

All polymers studied were polyesters, and dielectric measurements were limited to the solid state. I 02-1 07 Despite similarity in the dipole structure of these chains, they possess the same principal permanent dipole moments of the ester carbonyl groups. It is difficult to draw useful conclusions from these investigations. We therefore restrict ourselves to a summary of the main results obtained for different polymers.

Most dielectric studies have concentrated on copolymers obtained by different modifications of the PET chain.l03 IOS,107. Takase et al. l03

compared the dynamics of random copolyesters: a semi rigid hydroxybenzoic acid (60 mol%)/ethylene terephthalate(40 mol%) (HBA60/ PET 40) with a rigid hydroxybenzoic acid(75 mol%)/hydroxynaphthoic acid(25 mol%) (HBA75/HNA2S)' It is characteristic of these polymers that after solidification they exist in the liquid-crystalline glassy state, i.e. the orientational ordering is retained. DR measurements were performed in the frequency range 1 kHz-l MHz and a wide temperature range in the solid glass. Both copolymers exhibited a similar low-temperature ( < O°C, 1 kHz) f)-relaxation characterized by a small activation energy of the order of 20 kJ mol- I, probably associated with local motions of the polar ester groups. Apparently similar a-relaxation showed up in both polymers at higher temperatures (~75°C, 1 kHz), although the origins of these relaxations are probably completely different. For the stiff-chain HBA7S /HNA 2S , the a-process is characterized by activation energy of the order of 160 kJ mol- I, and probably involves the cooperative rotation of the neighboring segments on the same chain about the chain contour. The HBA60 /PET 40 chain is significantly more flexible because of the ethylene linkage. This enables the rigid segments to perform cooperative motion on a significantly larger scale, which is suggested by a large activation energy of about 500 kJ mol- I. Finally, at much higher temperatures a process associated with melting is observed: for HBA60 /PET 40 between 160°C and 180°C, and for HBA 75 /HNA 25 at lOODC higher. This large temperature discrepancy in the onset of cooperative interchain motion is attributed by Takase et al. 103 to the presence of the ethylene linkages in HBA60 /PET 40 and their absence in HBA75 /HNA 25 .

Boyd and his co-workers extended their earlier studies of the PET homopolymer l08 on HBA/PET copolyesters, being primarily interested


in the influence of the HBAx/PETy copolymer composition on dielectric relaxation. 104 Two random copolymers were studied: (x, y) = (60,40) and (80,20). In comparison with measurements of Takase et al.,103 the frequency range was expanded considerably downwards (5 Hz-50 kHz), which allowed them to observe in HBA60 /PET 40 a split of the (X-process into two (see Fig. 4.17). (Takase et al. in fact overlooked an indication of such a splitting, cf. figure 1 of Ref. 103.) Boyd and coworkers were able to add more significant information to that obtained by Takase et al. Firstly, the {J-process observed for HBA/PET copolymers shows a close resemblance to the sub-glass transition of the PET homopolymer with regard to activation energy and width. However, the temperature at which it appears rises with increasing content of HBA. The strength of the DR process appears to be independent of composition, (see Fig. 4.18). Altogether, these observations indicate that the {J-process should indeed be assigned to the ester group, as suggested by Takase et al.

In the complex (X-process, its overall strength depends linearly on the content of PET in the copolymer (see Fig. 4.18). It also occurs at significantly lower temperatures in copolyesters than in the pure PET. The split of the (X-process into two sub-processes in HBA60 /PET 40, Boyd and coworkers attributed to different PET-rich regions in the sample,

0-3,--------------------------------------,

w

200

Fig. 4.17. Temperature dependence of dielectric loss spectrum as a function of frequencies of (a) 5 Hz, (b) 50 Hz, (c) 500 Hz, (d) 5 kHz, (e) 50 kHz, for slowly cooled PET4o HBA6o ' (Taken with permission from Macromolecules, 20,988, ©

1987, American Chemical Society, Ref. 104.)


15

p 10 • ---

05

w 0 <J

',5

Ct

',0

0

0·5 ~/ /

0

¢pET

Fig. 4.18. Dielectric increment ~e' of the l/. and fJ processes as a function of the molar fraction $PET of PET in PETxHBA 1 - x copolymer. (Taken with permission from Macromolecules, 20, 988,:g 1987, American Chemical Society,

Ref. 104.)

although it is not a conclusive assignment. The most interesting result concerning the (X-process is a decrease of the glass transition temperature on increasing the content of the rigid HBA units, an effect which is yet to be explained.

Recently Kresse and colleagues reported preliminary results of DR studies on yet another type of LCP derived from PET. 107 They 'diluted' the ethylene terephthalate units in a chain by adding (in 1: 1 molar ratio) n-pentane-l,5-diyl-4-oxybenzoic acid (A) and 1,4-bis(4-acetoxybenzoyloxy)benzene (B). As a result, random copolymers PET xAY!2 BY/2 were obtained, which for x::; 80% form 9 nematic phase on melting. Measurements were performed for five different compositions of the copolymer, (x, y) = (20,80), (40,60), (60,40), (70,30) and (80,20) per cent, and for pure PET for comparison. The DR measurements were carried out in the very broad frequency range 1 Hz-l GHz, and over a


broad temperature range. As for other PET-related copolymers, at low temperatures the a and {J processes are observed for all samples in low-frequency range (f;( 10 MHz), and merge into a single high-frequency broad relaxation process at high temperatures (see Fig. 4.19). Both processes show dependence on the concentration of A and B components in the chain. Higher content of A and B units clearly leads to a substantial decrease of the strength of the a-process and an increase in its width. At the same time, the relaxation rate shows some tendency to increase slightly. These effects indicate changes in the micro-Brownian dynamics of the chain, although no explanation of the nature of these changes is available at the moment.

The {J-process is much more sensitive to the A and B unit content of the chain. The process strength increases significantly on increasing the fraction of the modifiers, i.e. it nearly triples on going from pure PET to PET 60A20B20' The relaxation rate shows a systematic behavior: on introduction of a small amount of modifiers the rate increases, but for x;( 80% the rate begins to increase linearly with concentration of

N I "-

<.>

7

6

4

3

25 3.0 3.5 4.0

IOOOK/T

Fig. 4.19. Temperature dependence ofthe relaxation rate fe of PET xA(l - x)/2 B(l - x)/2

copolymer as a function of the PET fraction present in the copolymer: (0) 100%; (+) 80%; (.~) 70%; (0) 60%; (e) 40%; (~) 20%. (From Ref. 107, with permission

of Springer-Verlag.)


modifiers (see Fig. 4.19). Interestingly, the process breadth seems to be essentially insensitive to variation of the chemical composition of the chain, as does the activation energy (57 kJ mol- l for PET, and 55-59 kJ mol- 1 for PET XAY/2 B y/ 2 ). Altogether, these observations confirm that local reorientations of the -COO- and C-O-C dipoles are responsible for the [J-relaxation. The increase of the process amplitude results from an increased number of dipoles taking part in the relaxation on introduction of modifiers. At the same time, it seems that for PET content below 80% the formation of extended aromatic blocks takes place, increasing steric hindrance to dipole reorientations and thus leading to a slowdown of the relaxation rate.

The C( and [J processes in PET xAY/2 BY/2 copolymers merge at high temperatures into a single C([J-relaxation which is detectable at high frequencies (see Fig. 4.19). The merging takes place over a very wide temperature range in which separation of the constituent C(- and [Jrelaxations is impossible. The C([J-process emerging at high temperatures shows Arrhenius behavior; it cannot, however, be considered a simple continuation of the glassy relaxation temperature behavior. Data of Kresse et al. lo7 suggest that bothC(- and [J-relaxation processes undergo significant changes in the temperature transition region, both successively losing their distinctive features and finally merging into one dynamical process.

The role of stiff segments in the [J-relaxation of the mcLCP was studied by Blundell and Buckingham. lo2 In particular, they investigated the role of 2,6-naphthalene rings in phenYl/naphthyl copolyesters. A study was made of three closely related polyesters which contain the same proportions of naphthalene and phenyl groups, but differ in the way in which the bordering ester linkages are joined to the 2,6-naphthalene groups. In each case different dipole moments were associated with the rotating naphthalene unit (see Fig. 4.20). Note, that, since it may be assumed that the carbonyl oxygen is coplanar with the neighboring ringlOg so that the carbonyl group and ring rotate as one entity, in the case of polyesters II and III in Fig. 20 there should not be an effective dipole moment rotating with the naphthalene moiety. DR measurements were performed on solid samples in the frequency range 50-10 kHz. Contour maps for dielectric loss of polyesters I and III are shown in Fig. 4.21. The most characteristic result of these studies is the virtual absence in II and III of the [J-process which is present in I. This result is important because dynamic mechanical analysis data of the same samples showed the presence of the [J-process in all three polymers. This leads to the


o II c \o-j-

(b)

o

-f &'~-o \oj f ,- \c-- II

o (c)

Fig. 4.20. Schematic representation of possible naphthalene moiety rotations in different mcLCPs: (a) polymer I, (b) polymer II, and (c) polymer III.

200 a 0.3 0.2

b~~; T(OC)

150

100

50

0 ~::; -50

2 3 4 2 3 4 log (fe/Hz)

Fig. 4.21. Temperature-frequency contour map of the dielectric loss spectrum for (a) polymer I and (b) polymer III. (From Ref. 102 by permission of the

publishers, Butterworth-Heinemann Ltd. [j)

204 Jozef K. Moscicki

conclusion that the {3-process observed mechanically must be associated with the naphthalene moiety.

We have to note at this point that despite the observed tendency of the polyesters studied to form the liquid-crystalline glassy state, no effort has been made so far, to study oriented samples of solid mcLCPs. As measurements of Davis and Ward have shown in the past, 110 such studies may provide us with more significant information on different mechanisms governing the a- and {3-processes than can be extracted from measurements on unoriented materials.

Finally, we mention the very interesting pioneering work of Vallerien et ai.,106 who studied a completely new LCP, namely the discotic main chain LCP, in which the mesogenic groups incorporated into the main chain are disc-like (see Fig. 4.22). It is characteristic of the polymer studied that the disc-like moieties were attached to the alkyl spacers via ester linkages, each possessing a significant dipole moment. DR studies performed in the frequency range from 10- 1 to 107 Hz and temperature range from - 170°C to 180°C revealed the existence in the glassy state of only a single broad asymmetric relaxation process, which can be consistently described by the empirical Havriliak-Negami and/or KohlrauschWilliams-Watt functions (cf. Section 4.2). The process was Arrhenius-like with activation energy of about 21 kJ mol-I. With the aid of 2H NMR data III it was deduced that the observed DR process was due to the hindered rotations of the ester groups, (see Fig. 4.22). What is interesting is that it was suggested that the ester groups were oriented perpendicularly to the neighboring aromatic rings, contrary to situation observed in copolyesters containing the rod-like segments (see above).

RO

Mn= 19100

OR

OR

OR

2H

2H 0 II

0/"'0-- CH 1-(CH 1in

Jm

9 330K 0416; R=-CsH" n=B

Fig. 4.22. Schematic representation of a disco tic polymer liquid crystal investigated by Vallerien et al. Arrows indicate rotations of the ester groups responsible for dielectric relaxation spectrum of this material. (From Ref. 106, with

permission of Steinkopff Verlag.)


4.5.2.2 Side chain liquid crystal polymers As discussed in Chapter 1 by Brostow, side chain liquid crystal polymers can be realized in many distinctive architectures. Dielectric relaxation studies were, however, performed mostly on a type of scLCP in which mesogenic groups are linked at one end to a backbone-chain via a spacer.18.24.29.112-162 In particular, these were the (i) polyacrylate-based and (ii) polysiloxane-based polymers.

Ra CH3

I I (i) -fCHz-C-+" (ii) -fSi-O-+m (4.48)

I I COOR R

synthesized by leading groups in the field [Shibaev and coworkers in Moscow, Ringsdorf and coworkers in Mainz, Finkelmann and colleagues in Freiburg, and Gray and coworkers in Hull]. Rare mesogenic groups, and the Ra group is H, CH 3 or Cl.

Dipole moments which can give rise to the dielectric relaxation signal are present on the main chain and on the side mesogen groups. The backbone motions can be dielectrically visible because of the ester group dipoles and of the Si-O bond dipoles present in polyacrylates and polysiloxanes, respectively. Note, however, that while the Si-O dipoles are rigidly attached to the chain backbone, the COO acrylic ester groups are adjacent to the backbone and, thus, take part not only in the motion of the main chain but also of the side chain. Dipole moments of the mesogen side unit are usually associated with the rigid core and with the terminal groups, one of which is the spacer in scLCPs.

DR measurements of scLCPs were initiated by Kresse and colleagues, who studied polyacrylate-based polymers/ 12 115 and most by far molecular dynamics studies of scLCP by means of DR spectroscopy were caried on these materials.116-132 Dielectric investigations of polysiloxanebased scLCPs were dominated by the research of Attard, Williams and colleagues, who studied extensively orientational effects in liquid-crystalline phases of scLCPs caused by a.c. electric fields.18.z9.133 1 S6

Most of the polyacrylate-based scLPCs investigated so far by means of DR spectroscopy are listed in Table 4.3. For convenience in the following discussion, we introduce in Table 4.3 a unified shorthand notation for different chains. Codes PA, PMA, and PCA correspond respectively to poly acrylate (Rl =H), polymethacrylate (Rl =CH3), and polychloroacrylate (Rl = Cl) backbone chains. The main chain code is followed by,


Table 4.3 Polyacrylate-based side chain liquid crystals investigated by dielectric relaxation

spectroscopy

rh CHz I ~O

Rl-C-C", 4-J 0-(CHz)n-RZ-C6 H4 - R3 -C6 H4 - R4

Rl Rz R3 R4 n Acronym Reference ----_.

H 0 CN 0 PA/O/-/CN 126 H 0 CN 2 PA/2/-/CN 126 H 0 CN 3 PA/3/-/CN 126 H 0 CN 4 PA/4/-/CN 126,130 H 0 CN 5 PA/5/-ICN 112,115,126 H 0 CN 6 PA/6/-/CN 122,126,129,130 CH3 0 CN 5 PMA/5/-/CN 113 H CO CN 0 126 H CO CN 5 126 H COO CN 0 126 H 0 COO CN 2 PA/2/COO/CN 116,118 H 0 COO CN 6 PA/6/COO/CN 116, 117, 119, 130 H/CH3 0 COO CN 6 123, 131 CH3 0 COO CN 6 PMA/6/COO/CN 123 H 0 COO OCH3 2 PA/2/COO/OCH3 117,118,127 H 0 COO OCH3 6 PA/2/COO/OCH3 117,118,120 CH3 0 COO OCH3 2 PM A/2/COO/OCH 3 117,118 CH3 0 COO OCH3 6 PMA/6/COO/OCH3 117,118,120 CH3 0 COO OC4H 9 2 PMA/2/COO/OC4H9 117,118,120 CH3 0 COO OC4H 9 6 PMA/6/COO/OC4H9 117,118,120,127 C1 0 COO OC4H9 2 PCAj2/COOjOC4H9 118 C1 0 COO OC4H 9 6 PCA/6/COOjC04H 9 117,120,121

separated by obliques, the length n of the spacer (number of (CH 2 ) groups in the spacer); the type of the interphenyl group, R3 ; and the type of the terminal group, R4 . The degree of polymerization, x, varied from a polymer to polymer, but Kresse et al. 113 found that this parameter has little influence on the molecular mobility and, consequently, it has been omitted in Table 4.3.

Not only is the number of polysiloxanes studied significantly smaller but they are also more difficult to classify in a similar manner to polyacrylates. Nevertheless, in Table 4.4 we have attempted to classify these materials, hoping it will aid dicussion.

Tab

le 4

.4

Pol

ysil

oxan

e-ba

sed

side

cha

in l

iqui

d cr

ysta

ls i

nves

tiga

ted

by d

iele

ctri

c re

laxa

tion

spe

ctro

scop

y

r-h

o I

CH

3-S

i-(C

H2

ln-R

l-C

6H

4-R

2-C

6H

3R

3,-

R3

: S

iO/n

/R2/

R3;

R3

4-J

Rl

R2

R3,R

3 n

Acr

onym

0 C

OO

C

H3

CN

5

SiO

/5/C

OO

/CH

3'C

N

0 C

OO

C

H3

CN

6

SiO

/6/C

OO

/CH

3'C

N

0 C

OO

C

H3

CN

8

SiO

/8/C

OO

/CH

3'C

N

CO

O

CN

4

SiO

/4/-

j-C

N

CO

O

CN

10

S

iO/l

O/-

j-C

N

0 C

OO

C

N

n [SiO

/3/C

OO

j-C

N]7

s/[

SiO

/6/C

OO

j-O

CH

3]92

5 0

CO

O

-O

CH

3 0

CO

O

CN

~}

[S

iO/3

/CO

Oj-

Cl]

75

/[S

iO/6

/CO

Oj-

OC

H3]9

25

0 C

OO

-

OC

H3

CO

O

CN

D

[SiO/7

/-j-

CN

]2s/

[SiO

/3/C

OO

j-O

CH

3]75

0

CO

O

-O

CH

3 0

CO

O

CH

3 C

3H

7 :}

[S

iO/4

/CO

O/C

H3·

C 3H

7] 5

0/[S

iO/6

/-j-

CN

] 50

0 C

N

Ref

eren

ce

137,

142,

149

14

3-14

5,14

9,15

2,15

4 13

8-14

1,14

6,14

7,14

9 15

1 15

1,15

3

134-

136

134-

136

155

148,

150,

152

o ii;'

;;;-~

.... ,,' :>0 ~ '" >< ~ o· '" ;;l" s:: '" " <S g ~ " IS" .... t-o

z' " is:. ()

.... ';;; ;::; 0;-

N o -..J


First to observe and classify different dielectric relaxtion processes present in scLCPs were Zen tel and coworkers, 11 7,118 who studied unoriented samples of several different poly acrylate-based scLCPs over a wide range of temperature, covering the isotropic, mesomorphic, and solid states (Fig. 4.23). They observed up to three high-frequency relaxations (namely, the f31, f3z, and y-processes) in the glassy state, and associated them with motions of dipoles on the side chain. Above the static glass transition temperature Tg , two additional processes were detected, the higher-frequency one showing typicaill-relaxation behavior and associated with the segmental motion of the main chain, and the low-frequency one (J-process) active in the mesomorphic and isotropic states and assigned to the end-over-end reorientation of the meso genic unit.

N

I

7

~5 OJ Q

3

3 4 5 6 7

1000 KIT

Fig. 4.23. Summary representation of the temperature dependence of the relaxation rate for five different dielectric relaxation processes observed in side chain polyacrylate liquid crystals. Tg denotes the glassy transition, and J, N, and S stand for the isotropic, nematic and smectic phases, respectively. The broken line

represents results of NMR investigations (see text).

Not all five relaxation processes are always simultaneously observed in individual poly acrylate-based materials. In fact, this finding significantly aids detailed assignment of the relaxation processes to particular modes of motion. By changing fragments of molecular architecture of the side chain unit or of the main chain and performing comparative DR studies it is, in principle, possible to identify different relaxation processes with particular dipole moments and their motions, and to study the influence of

Dielectric Relaxation in Macromolecular Liquid erntals 209

different functional groups on the intramolecular dynamics of the polymer chain.

Motions of the side chain dipoles in the glassy state. The most rapid DR process in the glassy state is associated with the reorientation of the mesogenic unit end-group dipole. It has been observed for polymers with an n-butyloxy end-group (PMA/2/COO/OC4H9, PCA/2/COO/OC4H9, PCA/6/COO/OC4H 9 ,117.118 PMA/6/COO/OC4H9,117.118.127) and

Zentel et al. 11 7.118 termed it the y-process (see Fig. 4.23). For example, for PMA/2/COOjOC4H 9, the process showed up at frequencies above 104 Hz at temperatures as low as 140 K (270 degrees below the static glass transition) rising to 107 Hz above 185 K.117 The temperature dependence of the process was Arrhenius-like with the activation energy ranging from about 17±4kJmol- 1 for PMA/6/COOjOC4H9127 to 24±4kJmol- 1 for PMAj2/COOjOC4H9.117 However, the same process was not detected by Zentel et al. for the methyloxy end-group, but Parneix et al. 126 reported a high-frequency process for PAj6jCOO/OCH 3 , which they assigned to the rotation of the terminal methyloxy group. However, high temperatures (not far below the static glassy transition) and the frequency range (~10 7 Hz) at which this process showed up, together with the fact that Parneix et al. 126 did not observe an additional relaxation process associated with the central ester group COO (see below), suggest an alternative interpretation of their results as being associated with the ester group rather than with the methyloxy one. Confusion is compounded by the fact, that the phase transition temperatures of the sample of Parneix et al. are in substantial disagreement with the temperatures quoted for PA/6/COO/OCH 3 by Zen tel et al.117.118 and Vallerien et al. 127

The PI-process was assigned by Zentel et al.117.118 to the rotation of the dipole moment of the central phenylbenzoate ester group (COO) (see Fig. 4.24). Because of the frequency range in which the PI-process shows up, we know that it survives the glassy transition unchanged; thus in this aspect it resembles the p-process observed in ordinary polymers (cf. Fig. 4.11). For PMA/2/COO/OC4H9 the PI-relaxation is five orders of magnitude slower than the y-process at 200 K (extrapolated) and has an Arrhenius-like temperature behavior, with activation energy of the order of 50 ± 7 kJ mol- I. II 7 The process appears for all kinds of mesogenic units containing the central ester group, independently of the spacer length and the end group, and with very similar values of the activation energy. II 7.118 The existence of this process was confirmed by Vallerien et al. 127 and Parneix et al. 126 (see above). The assignment of


+ --f Relaxation

n = 2, ... ,6 y Relaxation

R =H,CH3,Cl R2 = -OCH3, -OC4Hg, -CN

Fig. 4.24. Schematic representation of different molecular motions involved in the dielectric relaxation of polyacrylic liquid crystals.

the /31 DR-process to this particular mode of motion is strongly supported by 2H-NMR studies of the terminal phenyl ring of deuterated PA/2/COO/OCH3 and PA/6/COO/OCH3 [U. Pschorn, H.W. Spiess, B. Hisgen and H. Ringsdorf, Makrornol. Chern. (1986) 187, 2711]. It was found that in the liquid-crystalline glassy state the phenylene terminal rings undergo 1800 jumps about their local C raxes (~ the long axis of the mesogenic unit). The timescale of this process was identical to that observed for the /3cprocess. Also the mean correlation time dependence on temperature coincides with the temperature dependence of the /31-process; see the thin broken line in Fig. 4.23. These results strongly point to the reorientation of the central rigid core of the meso genic unit as a whole about the long axis in the glassy state. 127

Up to this point there was no apparent confusion with the assignment of DR processes observed in the glassy state of LCPs. However, additional relaxations are detected, interpretation of which is not as obvious as in the former two cases. For example, Zentel et al.117.118 reported the existence of an additional relaxation process, the /3rprocess, for PA/6/COO/CN (see Fig. 4.23), On the frequency-temperature plot it showed up as an intermediate between the y- and /3I-processes; it was also Arrhenius-like in temperature behavior with an activation energy of 35 ± 10 kJ mol- 1. Because this process was not observed for a nearly identical polyacrylate chain that had a shorter spacer, i.e. PA/2/COO/CN, Zen tel et al. 117 ,118 associated this process with motions within the spacer group. A similar DR process was also reported for PA/6/COO/OCH3 / 17 a polymer with a spacer of the same length (n=6) but with the mesogenic unit terminated by the methyloxy group instead of CN, but not for PMA/6/COO/OC4 H 9 ,II7.I27 which has the same


spacer, an n-butoxy terminal group, and a methacrylate instead of acrylate backbone. This might suggest involvement of the backbone flexibility in the process via the acryl ester dipole (the polymethacrylate chain is more rigid than the polyacrylate chain; see Section 4.4). To compound the confusion, Parneix et al. in their study ofPA/6/-/CN,127 a polymer with a mesogen unit without the central ester group, and for which, therefore, neither PI-relaxation nor y-relaxation should appear, observed two very weak relaxation processes active in the glassy state, one in the kilohertz and one in the megahertz frequency range. Unfortunately, Ie in both cases was determined with very poor accuracy, so the estimation of activation energy is unreliable. Nevertheless, in the absence of the ester central group and of the terminal n-alkyloxy group relaxations, it is reasonable to associate these processes with the spacer group dynamics and with libration of the terminal cyano group, respectively. 126 If so, the spacer group dynamics in PA/6/-/CN showed up in the same frequency range as the ester central group reorientation in PA/6/COO/OCH 3 , and the libration of the cyano group in the former occurred in the frequency range characteristic for, presumably, the spacer group dynamics in the latter! This really would be a surprising finding, which can be explained only if one assumes that the P I-relaxation is due to the dynamics of the whole side chain about the long axis of the meso genic unit, including the spacer with the acryl ester main chain group*, and that the Prrelaxation is caused by the libration of the longitudinal component of the side chain total dipole moment. The latter, is supported by a finding of Zen tel et al. that the P2-process is not observed for a short spacer (n = 2), since then the mesogenic unit is not sufficiently decoupled from the backbone for the libration to take place. Because of the very poor quality of the dynamical data of Parneix et al. for the glassy state,126 this interesting hypothesis requires further study.

Very sparse dielectric studies of polysiloxane scLCPs below the glassy transition show evidence of only one relaxation process, similar to the P I-process in polyacrylates, with a very broad distribution of relaxation times.136.156 The temperature dependence is Arrhenius and the activation energy of 51 kJ mol- I matches that of the PI-process in polyacrylates (see above). Measurements performed on the oriented smectic glass show that

*Becauseof the broad distribution of relaxation times characteristic of the !iI-process, it is probably safer to consider it as a superposition of relaxation processes resulting from reorientations about the long axis of different polar groups on the whole side chain.

212 }oze{ K. Moscicki

the mean relaxation time is independent of the sample orientation with respect to the probing electric field. 136 Pranoto et al. interpreted the process as probably associated with the motion of Si-O backbone dipoles. 136 In our opinion, however, the independence of the mean relaxation time from the sample orientation alone does not exclude participation of side-chain dipole moment reorientation about the long mesogen axis in this process. More probably, since only one broad relaxation process is observed, many different motions contribute to this relaxation process, involving the backbone and the side groups.

Dynamics of the mesogenic unit and of the backbone in the mesomorphic and isotropic states. Above the static glass transition two new intramolecular motions become active in side-chain liquid crystal polymers. Firstly, as in flexible-chain polymers, the micro-Brownian segmental motions of the backbone chain are successively liberated. The increased mobility of the backbone chain enables in turn large-scale reorientations of the long axis of the pendant mesogenic units. As a result, new DR relaxation processes appear.112-148

Dielectric relaxation measurements of the mesomorphic phases of scLCPs were initially carried out on unoriented samples.112.113.115 120 Kresse, Shibaev and colleagues were first to report on DR in the isotropic and liquid-crystalline (unoriented) phases of PA/5/-/CW 12 and PMA/5//CN. 113 The studied sample of the former is nematic below 350 K and exhibits the static glass transition at about 323 K. The latter is in the smectic-like state below 397 K, with the glass transition at 333 K. For both materials the DR spectrum is dominated by a strong relaxation domain with the frequency of the dielectric loss maximum fe in the range of 10-104 Hz in the nematic state, and 104_105 Hz in the isotropic state. The domain is well described by the Cole-Cole phenomenological equation (cf. Section 4.2, eqn 4.13)). The distribution of the relaxation times is larger for PMA/5/-/CN (the Cole-Cole parameter ex~0'33) than for PA/5/-/CN (ex ~ 0'22), but it does not change significantly at the clearing point, which suggests that the relaxation mechanism is the same in the isotropic and mesomorphic phases. The temperature dependence of the relaxation frequency fcC = 1/2nTo, cf. eqn (4.13)J for both materials is Arrhenius-like, with activation energy in the mesomorphic state twice that in the isotropic phase, see Fig. 4.25. (However, for unoriented polysiloxane SiO/5/COO/CH3 'CN, Attard et al. reported activation energy in the isotropic phase only 0·75 of that in the nematic phase, and a slightly non-Arrhenius temperature dependence of relaxation time. 137)

N I

u -en g


5~TSI I I I I

3 I

2.4

~T" I

SMECTIC I NEMATIC

R=H

2.6 2.8 3.0 1000 K / T

213

Fig. 4.25. Arrhenius plots for the is-process in the isotropic and unoriented nematic and smectic phases of homopoly acrylate and methacrylate with the same

side mesogenic unit.

Because the dipole moment of the cyano group (4'05 D) is several times larger than any other dipole moments present in the chains, the observed relaxation was associated with the reorientation of the meso genic unit about the backbone chain.

This process was subsequently observed and studied in several other polyacrylatel17-131 and polysiloxane134155 materials, and Zentel et al. named it the b_process. 117

Zen tel et al. were also first to observe another high-frequency DR process in an unoriented mesomorphic sample. 117 It appears at frequencies above the b-process and shows a very broad absorption curve. The temperature dependence demonstrates non-Arrhenian behavior, i.e. an increase in the activation energy with decreasing temperature. It was assigned to the micro-Brownian segmental dynamics of the backbone, i.e. the dynamic glass transition process, and termed the IX-process. 11 7

Although measurements performed on unoriented samples gave evidence of new intramolecular motions in melt phases of scLCPs, for studies of molecular dynamics in liquid crystalline phases DR


measurements on oriented samples are indispensable, i.e. those with the measuring electric field parallel (homeotropic alignment) or perpendicular (homogeneous or planar alignment) to the direction of alignment (cf. Section 4.3). On the one hand, when the probing electric field is parallel to the director, Ell 0, the end-over-end reorientation of the mesogen molecule becomes well isolated from the other, faster modes of motion: cf. eqns (4.42) and (4.43) and Fig. 4.7. On the other hand, the perpendicular geometry (E1.D) effectively eliminates from observation the end-over-end mode of motion, thus enhancing the visibility of the fast rotations about the mesogen long axis and librations of the axis. In monomer liquid crystals eqns (4.42) and (4.43) have proved to work quite well, explaining major features of the DR spectra. In scLCPs, however, additional effects may complicate interpretation of dielectric data.

First of all, dipole moments of polar groups on the backbone chain also contribute to the DR signal. Therefore, following the model considerations given in Section 4.3, one should expect, in general, to observe additional contributions from the backbone chain dipoles in dielectric spectra of mesomorphic materials (ef. eqns (4.42) and (4.43) and Ref. 131):

s{(w)-s (oo)=GI![(1+2Sm)uiF,1(w)

+ (1- Sm)u? PI(W)] + Gb91!(Sb).u~ pb(W) (4.49)

and

s!(w)-e1-(oo)=G~[(l-Sm)uiFl(w)

+ (1 + Sm/2)u~ Fi (w)] + Gb9 1- (Sb).u~ Fb(w) (4.50)

where G's are proportionality factors, gk(Sb), k= II or 1., is the backbone chain orientation factor, Sb is the backbone contour order parameter, and Fh(W) is the relaxation function of the backbone mode. Sm is the order parameter of the side groups and other parameters are as defined in eqns (4.42) and (4.43). In eqns (4.49) and (4.50) we assume explicitly that although the backbone chain contour may be ordered, but does not influence the chain dynamics, i.e. only the amplitude of contribution from the backbone dipole can be different in both principal geometries of the measurement.

We must also be aware that the dynamics of the meso genic unit in each of the angular coordinates, i.e. rx, fl, andy (ef. Fig. 4.4), can be significantly different, depending on the local environment, even on the small angular scale. For example, rotation about the long axis ("I) in scLCPs can be much faster than two small-scale motions of the long axis, i.e. precession


of the long axis about the director (O(), and fluctuation of the long axis about the director ({3) (see Fig. 4.7). Therefore, in contrast to monomer liquid crystals, where modes Fi(w) and F~I(w) appear in similar frequency ranges,2 (cf. Section 4.3), in the side chain LCP each of the modes can split into two distinctive contributions: one from the fast modes in the y coordinate, which would relax a part of f.1~, and one from slower reorientations of the long axis (0(, {3), which would relax the rest of f.1~. scLCPs offer an opportunity to investigate these effects by varying environment of the mesogen unit, e.g. by changing the chain architecture.

Additional potential phenomena we should be aware of, are possibly different rates of reorientation of different polar groups in the mesogen side-chain. This effect was sometimes observed even in monomer liquid crystals,2 but in polymer liquid crystals might be profound owing to much higher viscosity and coupling of the mesogenic unit to the backbone chain. In fact, relaxation data from the glassy state have already indicated that the meso genic polar groups can reorient independently of each other and with different activation energies. If this tendency survives into mesophases, the second term relaxation functions u~ r(w) in the r.h.s. of eqns (4.49) and (4.50) should be replaced by appropriate sums over contributions from constituent polar groups, Lu~FHw). Any cooperation in reorientation would result in some intermediate state.

As would be expected, by analogy to the monomer liquid crystals, dielectric relaxation measurements performed on aligned samples* revealed that a slow relaxation process is characteristic of the parallel geometry, while the dielectric relaxation observed perpendicularly to the director exhibits noticeably higher relaxation rates.121124.126.127. 130-132.138-157 Extensive studies by Attard et al. of unoriented, partially oriented, and homeotropically (E II n) or planarly (Kin) oriented samples of siloxanes demonstrated that the same processes are observed in unoriented and oriented samples, e.g. the mean relaxation times of both the 0(- and J-process show no significant change on going from an unoriented to an oriented sample,u8 151 The orientation, however, enhances the strength of each process.18.114.121-124.126.127 The (X-process

(Kin) is characterized by a very broad dielectric loss peak which consists of at least two overlapping relaxation regions; sometimes the peak clearly separates into two components, but usually it can be reproduced only by

*Magnetic fields were widely used to orient polyacrylates,121-124,126.127,130-132 while polysiloxanes were routinely oriented by ac electric fields,138-151


superposition of two phenomenological Fuoss-Kirkwood curves, a narrower low-frequency !Xl-component, and a weak and broader highfrequency !X2-component. The <5-10ss peak (Elln) on other hand, becomes significantly narrowed by the orientation, which enables observation of broadening at the high-frequency side. 18,29,114,121124,126,127.138151 In

this case also the spectrum can be satisfactorily decomposed into two Fuoss-Kirkwood phenomenological peaks, a strong and relatively narrow low-frequency b'-peak, and a broad and small-amplitude high frequency <5"-peak, the former retaining all essential characteristics of the overall <5-process. 18, 141,146,149, 151,155

The appearance of two different modes in each direction prompted Attard et al. to interpret the DR spectra of polysiloxanes in the mesomorphic phase with the aid of the rotational diffusion theory of Nordio et al. 141 (cf. eqns (4.42) and (4.43)). Although the theory of Nordio et al. is derived for monomer liquid crystals and assumes that molecules perform rotational diffusion in the presence of the nematic potential, Kozak et al. have recently shown that identical results can be obtained from the uniaxial symmetry of the liquid crystalline phase, without resorting to this assumption (cf. Ref. 18 and Section 4.3). Consequently, Attard 142 and Kozak 18 claim that they successfully identified the observed relaxation domains with the relaxation modes predicted by the theory. Accordingly, the <5'-process would correspond to the n(w) mode, i.e., the flip-flop reorientation of the long axis of the mesogen unit, and the b"-process to the FHw) mode, i.e., a combination of the rotation of the mesogen unit about the long axis, coupled with local fluctuations of the long axis direction about the local nematic director I8 ,29,142 (cf. Fig. 4.7). Fi and Fi modes were assigned to the !X1- and !X2-relaxations for Kin, although without specification of which of the processes belong to which of the modes. Their interpretation of the DR spectrum for Elln may at first sight be intuitively justified; however, the assignment for Kin is doubtful if not supported by additional studies.

In particular, they completely neglect the main chain Si-O dipole contribution to the relaxation spectrum. Kozak has found that the mean relaxation times for !xz- and and <5"-processes and the activation energies are nearly identicaj18 (see Fig. 4.26). This might be due to the fact that both Fi and pt modes involve similar small-angle reorientations of the long axis and the reorientations of the mesogenic unit about the smallest inertia axis. Thus they must be processes with similar reorientation rates (cf. Fig. 4.7), but on the basis of his data, it cannot be excluded that the high-frequency processes in the homeotropic and

N I

~ ..... Ol Q


71-

p- '-, 61-

I'-~-

51-

41-

33'10

'- ' "-

I 3·12

'. --. '-, '- '-, '-, '- '-"_ Fl '- "--~ -

I I I 3·14 3·16 3·18

1000 KiT

217

-

'''''''-. --

-

3,20

Fig. 4.26. Typical Arrhenius plots for four principal relaxation modes resolved from experimental data for polysiloxanes. (From Ref. 18.)

homogeneous geometries are also due to micro-Brownian motions of the backbone chains. Undoubtedly, more sophisticated experimental studies of scLCPs are needed to support the interpretation of Attard and his colleagues. Such experiments would, in particular, involve more detailed studies of the laterally attached side chains,132.133 since in this case the spacer restricts rotations of the mesogenic unit about the long axis, leaving the flip-flop motion essentially unperturbed. In fact, one might expect that the rotation about the long axis would become even slower than the flip-flop motion, in contrast to the dynamics of the end-attached mesogenic units.

Dielectric studies of polyacrylates are by far the most extensive. They involved a large number of different mesogenic groups and modifications to the main chain as welJ.116132 As a result, quite a detailed insight into the molecular dynamical origins of dielectric relaxation in the polyacrylates is possible. We begin with the perpendicular geometry of the dielectric experiment.

The (X-process observed in polyacrylates is most frequently characterized by a very broad dielectric loss peak, reflected in small values of the


Cole-Cole parameter; cf. Section 4.2, eqn (4.13). The distribution of relaxation times is dependent on the structure of the mesogenic unit, in particular, on the dipole moments present on the unit. For example, the distribution broadening increases on going from the cyanobiphenyl side group (PAj6j - JCN) to the cyanophenylbenzoatoxy side group (PAj6jCOOjCN)122,127 and to the methyloxyphenylbenzoatoxy side group (PAj6jCOOjOCH 3 ).126 This behavior clearly indicates that the (X-process does not only involve the dipoles of the ester group at the backbone-chain. Because the probing electric field is perpendicular to the average orientation of mesogenic groups, the transverse components of the dipole moments of mesogenic units also contribute to the dielectric signal. In the first case, the oxy-group dipole moment (R2 group in the molecular scheme in Table 4.3) is active; in the second, there is an additional polar interphenyl COO group (R3), and in the latter, the cyano group (R4) is replaced by the methyloxy group. To add to the complexity of the picture, a libration of the longitudinal component of the mesogenic group total dipole moment is probably also present in the spectrum, (see eqn (4.50)). Therefore, since in general mobilities of different polar groups can be different and exhibit different temperature behavior, the breadth of the resultant absorption spectrum reflects the diversity of dipoles and motions involved. Not surprisingly then, the temperature dependence of the position of the absorption maximum on the frequency scale (reciprocal of the mean relaxation time) is strongly non-linearly dependent, especially close to the clearing point. For example, Bormuth and Haase123 observed practically zero apparent activation energy for the (X-process in PMAj6jCOOjCN over a temperature range nearly 30 K below the clearing point (see Fig. 4.27). Similar retardation and temperature-dependent dielectric loss curve shapes of the (X-process on approaching the clearing point were observed for PA 122 and PCA,121 but in general these effects are markedly stronger for side groups with a strong positive dielectric anisotropy (111 ~ 111) than for those groups with a strong negative dielectric anisotropy (111 ~ 112)' To investigate the (X-process in more detail, Bormuth and Haase131 separated the spectrum for methacrylate-rich PA-coPMAj6jCOOjCN copolymers and PMAj6jCOOjCN homopolymer into two components, the (X1- and (Xrprocesses. The low frequency (Xl-process shows a narrower dielectric loss curve than that of the O(rprocess, and its strength increases on approaching the transition to an isotropic melt. They found that the separation of the (Xl-process decreases with increasing flexibility of the backbone chain, becoming inseparable for pure PAj6jCOOjCN. Also,


Xx x x x Xx

x x

-;:; 3 0: iso o : II :r: x . l

[lin

-1 L--2::-l.6::------':-'2. 7::L---,2~.8::----2;;-';.9::----:;"3.0;;----"-'

1000 K/ T

219

Fig. 4.27. Temperature dependence of the relaxation rate in the isotropic and oriented nematic and smectic phases of PMAj6jCOOjCN. (From Ref. 123, with

permission of Gordon and Breach Science Publishers, SA.)

the replacement of a mesogenic group with a strong longitudinal dipole moment by a group with a strong transverse dipole moment leads to the disappearance of this process. 136 Therefore, the ai-process is associated with the reorientation of the meso genic unit long axis. The arprocess would then be associated with rotations of the transverse component of the side chain dipole moments, but a tendency for the strength of the high-frequency arprocess to decrease with temperature and with stiffening of the main-chain is typical for the a-process in amorphous polymers, which suggests also a participation of the backbone ester polar groups in this process.

We note at this point that the geometry of the side group has a profound influence on the frequency range in which the overall a-process shows up. Pranoto et al. 136 investigated copolymers [SiOj3/COO/Rhs [SiO/6/COO/OCH3 ]92.s, where R=CN or Cl, and were not able to observe the a-process in the frequency range studied (10-107 Hz). However, Attard et al.,138-149 who investigated homopolysiloxanes SiO/n/COO/CH3 CN, i.e. with essentially the same mesogenic group but widened by the addition of the CH3 group to the second phenyl ring at the 2-position as indicated in Table 5.4, found this process only an order of magnitude faster than that of the (j-process. This finding supports our speculations above about possible significant decoupling of molecular rotation of the mesogen about the long axis, and local fluctuations of the long axis, in side chain polymers. It also suggests further DR experiments with various side group geometry


as a possible method for extracting a contribution to the overall c<-process from the backbone dipoles.

The J-process, observed in the parallel geometry of the oriented sample, is almost always characterized by a slightly non-Arrhenius temperature dependence, progressively so at low temperatures. The strength of the process shows a clear dependence on the longitudinal component of the mesogenic unit dipole moment. For example, at comparable temperatures, for PA/6/COO/CN, i.e. for the cyano terminal group with a large dipole moment parallel to the long axis, the dielectric loss maximum e;;' in the parallel geometry is of the order of 2, while for PA/6/COO/OCH3, i.e. for the methoxy end group with a ten times smaller longitudinal component, it is smaller by a factor of 2 or more. 126

Although, the typical relaxation frequencies are smaller by a factor of 103

and the activation energy is greater by a factor of 2, for PA-scLCP in comparison with the monomer analogues/ 26 all facts strongly support the proposed molecular interpretation of the J-process as the endover-end reorientation of the side group [F\(w) mode, cf. Eqns (5.42) and (5.49)]. However, in comparison with monomer liquid crystals, in scLCPs this process is a highly cooperative one. 129,153

The slowdown of the reorientation rate, increase of the activation energy, and the distribution of relaxation times in PA-scLCPs reflect the presence of the anchoring, the different environments and the dynamics that a mesogenic unit experiences in scLCPs in the isotropic and liquid crystalline states. This can be nicely demonstrated by variation of the polymer architecture.

Parneix et al. 126 elongated the spacer of PAln/-/CN from n = 2 to n = 6. Comparing the activation energy and the mean relaxation frequency for these materials in the isotropic state (no ordering potential), they found strong evidence for decoupling between the side unit and the backbone on elongation of the spacer (see Fig. 4.28). Although the decoupling saturation is not reachcd, it is clear from Fig. 4.28 that the mobility of the side chain increases rapidly with the elongation of the spacer, and for n;?: 6 the mesogenic unit decoupling from the main chain saturates.

It is well known for ordinary polyacrylics that the chain stiffness can be modified significantly by varying the alpha group.51 The influence of the flexibility of the main chain on the dynamics of the mesogen side chain was studied, on the one hand, by comparing the J-process for acrylate, methacrylate and chloroacrylate main-chains with the same spacer (n = 6) and very similar side groups,113.117,118,123 and on the other by varying the amount of methacrylate units on the PA-coPMA/6/ COO/CN


7.5 160

7.0

140

6.5 'j' roo.

N (5 I '-.. 120 ~ _u6.0

J "'-" ~ Ol '-.. 0 - 5.5 w

100

5.0

4.5 1 7 80

n

Fig. 4.28. The relaxation rate Ie and activation energy E as a function of the spacer length n for PA/n/-/eN. (From Ref. 126, with permission of Taylor and

Francis Ltd.)

copolymer. 131 The former approach seems justified, since the main chain substituent had very little if any effect on the relaxation strength. The most important findings from such a comparison are that the distribution of the relaxation times broadens and the activation energy increases in the sequence acrylate-methacrylate-chloroacrylate, i.e. as the segmental mobility of the backbone chain is increasingly restricted. Also in the case of PA-coPMA copolymer the width of the dielectric loss curve increases and shows signs of asymmetry with increasing methacrylate content (i.e. with chain stiffness).131 At the same time, the mean relaxation time becomes successively shifted towards lower frequencies as the activation energy increases. All these results must be considered clear and strong evidence of a significant degree of cooperation between the end-over-end reorientation of the side mesogenic group and conformational motions of the main chain.

Broadening of the distribution of relaxation times is also found to be caused by interactions between the mesogenic units. Haase et al. 122 and Parneix et aU 26 compared the dielectric properties of two polyacrylic scLCPs with slightly different meso genic units: PA/6/-/CN and PA/6/COO/CN. Both materials demonstrate a profound difference in the quasi-static dielectric permittivity and dielectric anisotropy: £; (PA/6/ -/CN)~12 versus £:(PA/6/COO/CH)~19, and (£;-G~)~6'5 for the former versus '" 12 for the latter126 Consequently, the strength of the


b-process is also different; in the nematic phase at temperature T~ TNdl·2 for PA/6/ - /CN the dielectric increment [e;I(<X))-e~(O)] is about 8, and for PA/6/COO/CN it is about 14.122 Such a striking change in dielectric parameters cannot be explained only by the presence of the interphenyl COO group dipole moment in PA/6/COO/CN and its absence in PA/6/-/CN. Most probably, relatively low values for PA/6/ -/CN are due to the presence of the antiferroelectric alignment of the cyano group dipole moments in this compound, while increased values in the case of PA/6/COO/CN can be explained by decorrelation of the dipole moments caused by the presence of ester interphenyl groups, since similar effects have been observed for monomer liquid crystals of a similar architecture. 164 In agreement with such a hypothesis is the distribution of relaxation times, which is larger for PA/6/-/CN than for PA/6/COO/CN,122 since decorrelated mesogenic units of the latter should move more independently than in the former.

What is interesting and obviously argues in favor of a high degree of cooperation between the main chain and the side chain in their dynamics is partial parallelism of the temperature dependence of the iY.- and b-processes, observed well within the mesomorphic state. 116,122,123, 125-127,131,141,146,150,154 Heinrich and Stoll studied the iY.- and b-pro-

cesses in PA/2/COO/CN and PA/6/COO/CN under high pressure,12S They found that the b-relaxation is closely correlated to the iY.-relaxation (see Fig. 4.29(a)), A most surprising finding is the pressure dependence of the mean relaxation frequency for the iY.- and b-processes. It turns out that not only is the IX-process behavior similar to that of amorphous polymers, but that both processes essentially parallel each other over the range of pressure studied (see Fig. 4,29(b)).165 This is clear evidence of a very strong dynamical coupling between the mesogenic units and the backbone chain, i.e. essentially cooperative dynamics. In fact, since both processes show significant non-Arrhenian temperature behavior, Attard et al. applied eqn (4.50), neglecting the dependence on the order parameter, to the b'- and iY.-processes in SiO/8/COO/CH3CN in order to estimate the temperature at which both motions cease to exist, and found that Tg-To(iY.)c:d6K, and Tg-To(b)c::::56K/ 46 which are typical values for glass-forming polymers.49 As can be seen in Fig. 4.29, both processes show also a tendency to merge at lower temperatures. 125 ,146 This phenomenon can be explained as originating from a dependence of both relaxation rates on the collective configurational entropy.166 At temperatures close to the glassy transition, the configurational entropy, a significant part of which is due to the polymer chain and spacer, becomes very


a

-2 2~.3--~2~.5~~2.7~~2~.9--~3.1~

1000 K/ T

(a)

:l\ ~·1 .l\J\ J

1000 3000 5000 pi bar

(b)

Fig. 4.29. Relaxation rate of the 0;- and fj-processes as a function of (a) temperature and (b) pressure for PAj6jCOOjCN (From Ref. 125, with permission of

Steinkopff Verlag.)

small, so we may expect that the relaxation rates of both processes would become similar in value and very small.

Finally, it has to be noted that the frequency of some reorientations can be moved to much higher frequencies, and dielectric spectroscopy of scLCP mesophases has so far not had such high-frequency (> 107 Hz) investigation. The only measurements performed up to 1 GHz are due to Parneix et al.,126 who used a wide frequency apparatus operating in the frequency range 5 Hz - 1 GHZ.126 Apparently, they isolated the methoxy terminal group relaxation in the nematic and smectic phases of PA/6/COO/OCH3 . The relaxation process shows up in both basic geometries of the experiment and in the frequency range 107-108 Hz, and exhibits Arrhenius-like temperature dependence. On going to the isotropic phase it moves out of the range of the apparatus. This study clearly demonstrates that much can be learned from high-frequency domain measurements in scLCPs.

4.5.2.3 Main chain/side chain liquid crystal polymers A separate class of liquid crystal polymers is polymers in which meso genic units are both incorporated into the backbone chain and attached as the side groups. These materials were synthesized in the course of searching for polymers which combine favourable properties of mcLCPs, such as desirable mechanical properties, broad range of thermodynamic stability of mesomorphic phases and exceptionally high orientational order, with the good orientability by external magnetic and electric fields and good optical, electro-optical and non-linear optical properties of scLCPs.


In a main chain/side chain LCP (mc/scLCP), the molecular dynamics, and the side chains dynamics in particular, reflect a new environment, i.e. higher orientational order and different restrictions arising from coupling of the mesogenic groups in side chains and in the main chain, for the mesogen units. Systematic dielectric relaxation studies have been performed on several mc/scLCPs, which are listed in Table 4.5.157-16°. Endres et al. studied numerous polymers with the side mesogen azobenzene unit terminated by the cyano-group with a strong longitudinal dipole moment, 15 7,159 while Kremer et ai, investigated mc/scLCPs with side groups the predominant component dipole moment of which is transverse to the long axis, 158, 156 X-ray studies of the former materials revealed that in the nematic and smectic phases aromatic rigid mesogen units tend to order parallel to each other, 159 as indicated schematically in Fig. 4.30. This mutual orientation of the backbone and of the side chain mesogen units has a profound influence on the dynamics of the latter.

Up to four different relaxation processes are observed in mc/scLCPs, although not for a single polymer (see Fig. 4.31). All investigated polymers display at least two relaxation domains, the high-temperature (low-frequency) and the low-temperature (high-frequency).157 156 Often the high-frequency domain is composed of two peaks, say, fJ' and fJ/I.157,15S Measurements performed on the oriented samples revealed that these processes are orientation independent and show surprising variation of their relative magnitudes with temperature; i.e. the low-temperature fJ/I peak increases in strength with temperature, while the hightemperature fJ' peak decreases at the same time (cf. figure 4 of Ref. 158, and figure 7 of Ref. 159). The depenence of relaxation frequency on temperature is Arrhenius-like (cf. Fig. 4.31), and the activation energy for the former is nearly twice that of the latter (64 kJ mol- 1 and 33 kJ mol- 1

for polymers studied by Endres et al./ 57 and 55-70kJmol- 1 and 30-40 kJ mol- 1 for those studied by Kremer et al. 15S Kremer et al. assigned both processes to rotational motions of the meso genic units about their long axes: the low-temperature f3" process to the rotation of the meso genic side group, and the high temperature fJ/I-process to the rotation of the main-chain mesogenic unit. However, they did not elaborate on the peculiar interdependence of magnitudes of both processes. Endres et al. proposed a similar interpretation but suggested that this interdependence may result from different local environments that each type of mesogenic units, the main chain and side chain, experience. 157 If this is so, the meso genic units, which would otherwise feature the same rotational dynamics about the long axis, are divided in less mobile and more mobile

Tab

le 4

.5

Mai

n c

hain

/sid

e ch

ain

liqu

id c

ryst

al p

olym

ers

inve

stig

ated

by

diel

ectr

ic r

elax

atio

n sp

ectr

osco

py

TY

PE

k [-l-

fH ~ A

-O-'CH

,k-O

0-,CH

,),-0-1

(C

H2)

6 I ,---

o

Ma

in c

hain

mes

ogen

0 B

r

Side

cha

in m

esog

en

-@-N

~ N

----@

--@

-N1N

--@

----

-@--

@-

-@-@

-P

olym

er

x y

Pol

ymer

x

y P

olym

er

x y

Pol

ymer

x

y

--@

-N=

N-@

-CN

1

6 6

3 2

2 4

6 6

2 6/

2 6/

2 C

HJ

-@

-N

=N

-@

-C

N

5 6/

2 6/

2

0

-@

-N

I N

-@-OCH2~HC2H<

39

6 6

40

6 6

37

6 6

CH,

-@-{

§)-

oC

H'1

HC,H

< 43

6

6 44

6

6 41

6

6

CH,

~OCH21HC'H<

31

6 6

32

6 6

29

6 6

Br

CH,

TY

PE

B: t o-@-co-o~o-

CO

---@

-O(C

H,)

, t p

oly

mer

(x,

R):

6(6

, O

CH

3),

7(9

, OC

H3

),

¥ m

8(

10, O

CH

3),

9(1

0, C

N)

(CH

,),

\ fn

\ .. ~

O--

-&--

N=

N--

-&--

R

I::l

1;;' " ~ .... ;:;'

:;.:, '" ;:;- >< ., 5' ;:

, s' ~ r, ci ;:; '" " " '" ;:;- .... l .E

;' '" E.: ~ ~ tv

tv

Ul


Nematic Smectic

n

Fig. 4.30. Schematic representation of mesogen unit order in the nematic and smectic phases of a main chain/side chain polymer liquid crystal. n is

the director.

6

I I let

I

\ \

\ \

\ \

\ \

\ \

\ \ /3"

OL-~ ____ ~~ ____ ~ ____ ~ 3 5 7

1000 IT

Fig. 4.31. Summary representation of the temperature dependence of the relaxation rate for four different dielectric relaxation processes observed in main

chain/side chain liquid crystal polymers (cf. Table 4.5).


fractions, with the ratio between the fractions varying with temperature. In favor of such interpretation is the presence of both processes in polymers 40 and 43 of see Table 4.5, which possess polar mesogenic groups only in the main chain (43) or only as the side chains (40).

Each of the polymers studied exhibits the high-temperature (lowfrequency) process active in the mesomorphic phase, but the origins and character of this process depend dramatically on the orientation of the side mesogenic unit dipole with respect to the long unit axis. For polymers with the cyano terminal group, i.e. with the strong longitudinal dipole moment, the observed relaxation is associated with the flip-flop of the unit about the short axis/ 57 .159 i.e. it is the same c5-process as observed in scLCPs. However, for polymers with the large transverse dipole on the side meso genic group, the hightemperature process has features of the IX-relaxation process, i.e. of the dynamic glass process.

The c5-process has characteristics of end-over-end reorientations of the long axis. It is characterized by having nearly the Debye-type of Cole-Cole plot, and its magnitude is similar to that observed for scLCPs. However, in contrast to monomer liquid crystals with the cyano end group and scLCPs with a similar side chain mesogen, where strong anti parallel coupling between the neighbouring dipoles is observed,23.126 the magnitude of the process in mc/scLCPs corresponds to the completely decoupled longitudinal components of the total dipole moment. 157 The frequency range in which the c5-process shows up depends visibly on the main chain stiffness. For example, at 100°C, for the polymer main chain with a flexible section of 12 methylene units (polymer 1 in Table 4.5), the relaxation frequency is about 5 kHz, but it decreases to 0·1 kHz for a chain containing only 4 methylene groups in the flexible section (polymer 3). In contrast to most of scLCPs, the temperature dependence of the relaxation frequency Ie is Arrhenius-like with activation energy in the range 115-140 kJ mol- 1. No significant correlation between the chain architecture and the activation energy is noted.157.159 However, the activation energy is surprisingly low in comparison with values observed for scLCPs with end-fixed mesogenic groups (150--200 kJ mol- 1, see above). This difference becomes even more pronounced for side-fixed mesogenic groups (> 240 kJ mol-1132.133). Clearly, coupling between the side group meso genic units and the backbone chain seems to be weaker for mc/scLCPs than for scLCPs. This can reasonably be explained as

228 laze! K. Moscicki

due to the strong mutual parallel alignment of the side and main chains, leading to much easier flip-flop reorientation of the side groups in mc/scLCPs than in scLCPs. In the latter, large changes in the backbone conformation are required in order to facilitate such reorientation. 159

Surprisingly, the a glassy transition process seems to be absent in polymers 1-9 in Table 4.5. It is observed, however, in polymers studied by Kremer et al., that, in turn, the b-process does not show up, probably owing to too small a longitudinal component of the mesogen dipole moment. 15S

The a-process is the dynamic glass process of the polymer backbone chain and has been observed for polymers featuring the glassy transition and crystallization. The temperature dependence of the mean relaxation frequency is strongly non-Arrhenian. Application of eqn (4.50) yields Tg - To (a) ~ 32 K, which is typical for glass-forming polymers and in good agreement with 36 K found for scLCPS. 146 The strength of the a-process was found to be largest for a polymer with 3-bromobiphenyl meso genic groups in the main chain and in the side chains (polymer 29), decreasing upon any substitution of these groups by other meso genic groups, and becoming weakest for polymer 39, which has only azobenzyl mesogenic groups in its structure. IS8 This result demonstrates that the side groups and the backbone meso genic units in the mc/scLCPs studied are coupled. Therefore, a spacer of 6 CH 2 units, which is sufficient to decouple the side mesogenic unit from the backbone in scLCPs (see above) does not fully decouple the glass relaxation of the backbone from motions of the side units.15s

For those polymers which undergo transition to the crystalline state, the a-process has strongly temperature dependent strength/57.15s probably owing to the crystallization process.

It has to be noted that Kremer et al. conducted their studies on purposely unoriented samples. 1s8 They did so because the presence of the optically active end groups in the side mesogen chains usually leads to chirality of the mesophase. As a result, polymers 31, 37, 39, 41 and 43 exhibit the cholesteric (chiral nematic) phase, and polymer 40, 41 and 43 the chiral smectic C* phase. 1s8 Since the smectic C* shows ferroelectricity, in order to separate the molecular rotational modes from the ferroelectric Goldstone and soft modes (cf. Section 4.3), samples should be unoriented. On the other hand, measurements performed on the oriented sample of the chiral C* phase of polymer 43 led Vallerien et al. to the observation of ferroelectric modes. 160


4.5.2.4 Ferroelectric liquid crystal polymers Ferroelectric side-chain liquid crystal polymers (fLCPs) have been synthesized only recently.156.158.167 (see Chapter 8). Because of the

dilution of chiral side chains in the backbone matrix, the spontaneous polarization is weaker than in the monomer chiral smectics, although Kapitza and others recently reported polysiloxane-based fLCPs with spontaneous polarization comparable to that observed for monomer liquid crystals. 168

Dielectric studies of the ferroelectric polymer and elastomer liquid crystals are only in their beginnings. 1 56.160 162 Therefore, we can only briefly summarize the most recent findings of the Mainz group. 156, 160.162 fLCPs synthesized by that group are routinely investigated by dielectric relaxation spectroscopy. The synthesized polymers were either of mc/scLCP type158.160 or of scLCP type/ 56 ,162 with various side chain

chiral groups employed. In the case of the mc/scLPC type materials, the backbone mesogenic unit was also varied. 158

The 3D plot of the dielectric loss £" versus log U/Hz) and temperature obtained for a thin ( '" 1 0 .urn) sample of polymer 43 (Table 4,5) oriented with the helix axis parallel to e1ectrodes l60 is shown in Fig. 4.32. Polymer

Soft mo(le IS~-phcsel

Soft mode IS;-S~ tronsi tlonl

I I I I II

\

Conductivity contribution

I

Fig. 4.32. 3D plot of the dielectric loss versus frequency and temperature for a thin oriented sample of the ferroelectric liquid crystal polymer 43 (Table 4.5).

(From Ref. 161 with permission of Hiithig Publishing Ltd.)


43 exhibits smectic A *, C* and J* phases, (cf. Table 4.5) and a weak spontaneous polarization. As the temperature is varied, the dielectric loss factor shows maxima differently located on the frequency scale, depending on the temperature range considered. The most striking feature, which can be seen in Fig. 4.32, is the lack of a visible temperature dependence of the positions and strengths of maxima within each phase, and the extraordinary strength of these maxima, being at least one order of magnitude stronger than the strength of the molecular rotational relaxation process. 158 Both features are characteristic of ferroelectric modes.

A strong, high-frequency ('" 100 kHz) peak which shows up in the vicinity of the A~C* transition is characteristic of the soft mode. It appears in a similar frequency range and has similar strength to the soft mode process in the monomer C* phase (cf. Fig. 4.9). On reducing the temperature, immediately below the phase transition the intensity of the soft mode becomes suppressed while the whole maximum is shifted to higher frequencies, also in agreement with results for the monomer C* phase (cf. Section 4.3.2). Both parameters, the strength and characteristic frequency, remain essentially constant within the C* phase. However, once the polymer undergoes the transition to the J* phase, the soft mode is restored to its position and strength from the vicinity of the A~C* transition. Within the J* phase, the frequency and the strength continue to be temperature independent.

The Goldstone mode appears once the polymer is cooled into the C* phase (cf. Fig. 4.32). By comparison with monomer chiralliquid crystals, the Goldstone mode frequency range ('" 1 kHz) is shifted by about one decade to lower frequencies. The intensity of the process depends on the magnitude of spontaneous polarization. For polymer 43, which has a small spontaneous polarization, the intensity of the process is very substantially reduced (compare Figs 4.9b and 4.32), but the Mainz group reported ferroelectric main chain siloxanes with very high spontaneous polarization, comparable with that observed for monomer chiral liquid crystals. 157 Consequently, for polymer A in Table 4.5, they observed an extraordinary strength of the Goldstone mode, similar in value to that typical of monomer materials.

The temperature dependence of the characteristic frequency of the Goldstone process in the C* phase depends on the material studied. For example, for polymer 43 it appears to be temperature insensitive, but at the transition to the J* phase the Goldstone mode becomes rapidly shifted to much lower frequencies. 16o In systems with the glassy transi-


tion, like polymer A, the frequency of this process is visibly temperature dependent, decreasing dramatically as the glassy transition is approached. 156 Similar behavior is observed for ferroelectric elastomer liquid crystals. 161

ACKNOWLEDGEMENTS

I express my special thanks to Professor Jack H. Freed for his hospitality during my visits to Cornell University. Professors W. Haase, H. Kresse, V. Shibaev, R. Zentel, and Drs C. Boeffel and G.R. Mitchell, supplied me with their recent publications. Permission to reproduce certain figures, or use figures partially, has been given by the editors of the following journals: Colloid & Polymer Science (Steinkopff Verlag, Darmstadt), Journal of Chemical Physics (American Institute of Physics, New York), Liquid Crystals (Taylor & Francis Ltd, London), Macromolecules (American Chemical Society), M akromolekulare Chemie (Hiithig & Wepf Verlag, Basel), Molecular Crystals and Liquid Crystals (Gordon & Breach Science Publishers SA, Montreux), Polymer (Butterworth-Heinemann Ltd., Guildford), and Polymer Bulletin (SpringerVerlag, Heidelberg).

This work was supported partially by the Polish Accademy of Science under project CPBP 01.12 and by NSF Grant DMR-88-18885-A02.

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F.E. Plenum Press, New York, 1972, p. 17; and references therein. 63. Miller, W.G., Annu. Rev. Phys. Chern, 1978,29,519. 64. Doi, M., J. Physique (Paris), 1975,36, 607. 65. Doi, M. & Edwards, E.F., J. Chern. Soc., Faraday II, 1978,74,560,918. 66. Doi, M., J. Polymer Sci. Polymer Phys. Ed., 1981, 19, 229; 1982,20, 1963. 67. Mori, Y., Ookubo, N., Hayakawa, R. & Wada, Y. J. Polymer Sci. Polymer

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642. 71. Bawden, F.C & Pierce, N.W., Proc. R. Soc. Lond. Ser. B, 1937, 123,274. 72. Freudlich, H., J. Chern. Phys., 1937,41, 1151. 73. Aharoni, S.M., J. Macromol. Sci. Phys., 1982,821, 105. 74. Onsager, L., Ann. N. Y. Acad. Sci., 1949,51, 627. 75. Flory, PJ., Proc. R. Soc. Lond. Ser, 1956, A234, 73. 76. Aharoni, S.M., J. Polymer Sci. Polymer Phys. Ed., 1982, 19, 281. 77. Aharoni, S.M., J. Polymer Sci. Polymer Phys. Ed., 1980, 18, 1303. 78. Roviello, A., & Sirigu, A., J. Polymer Sci. Polymer Lett. Ed., 1975, 13,455. 79. 1ackson, WJ. & Kuhfuss, H.F., J. Polymer Sci. Polymer Chern. Ed., 1976, 14,

2043. 80. Gullion, D. & Skoulios, A., Mol. Cryst. Liq. Cryst. 1978,49, 119. 81. Griffin, A.C & Havens, SJ., Mol. Cryst. Liq. Cryst., 1978,49,239. 82. Millaud, B., Thierry, A., Strazielle, C and Skoulios, A. Mol. Cryst. Liq.

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55. 89. Finkelmann, H., Ringsdorf, H. & Wendorff, lH., Makromol. Chern., 1978,

179, 273. 90. Shibaev, V.P. & Plate, N.A., Polymer Sci. USSR, 1978, A19, 1065. 91. Ringsdorf, H. & Schneller, A., Br. Polymer J., 1981, 13, 43. 92. Boeffel, c., Spiess, H.W, Hisgen, B., Ringsdorf, H., Ohm, H. & Kirste, R.G.,

Makromol. Chern., Rapid Commun., 1986,7, 777. 93. Boeffel, C. & Spiess, H.W., Macromolecules, 1988,21, 1626. 94. Miller, W., Wu, c.c. Wee, E.L., Santee, G.L., Rai, lH. & Goebel, K.G., Pure

Appl. Chern., 1974,38, 37. 95. Kulichikhin, V.G., Kudryavtsev, G.I. & Papkov, S.P., Int. J. Polymer

Mater., 1982,9,239. 96. Aharoni, S.M., J. Polymer Sci. Polymer Phys. Ed., 1980, 18, 1439. 97. Moscicki, lK., J. Polymer Sci. Polymer Phys. Ed., 1985, 23, 327. 98. Moscicki, lK., Rosato, V. & Williams, G. unpublished results. 99. Moscicki, lK., in Dynamical Processes in Condensed Matter (Adv. Chern.

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Nonlin. Opt., 1988, 157,455. 106. Vallerien, S.U., Kremer, F., Huser, B. & Spiess, H.W., Colloid Polymer Sci.,

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93. 108. Coburn, lC. & Boyd, R.H., Macromolecules, 1986, 19, 2238. 109. Hummel, lP. & Flory, lP., Macromolecules, 1980, 13, 479. 110. Davies, G.R. & Ward, I.M., J. Polymer Sci. Part A-2, 1972, 10, 1153. 111. Huser, B. & Spiess, H.W., Makromol. Chern. Rapid Commun., 1988,9, 337. 112. Kresse, H., Talroze, R.V., Makromol. Chern. Rapid Cummun., 1981,2, 369. 113. Kresse, H., Kostromin, S. & Shibaev, V.P., Makromol. Chern. Rapid Com-

mun., 1982, 3, 509. 114. Kresse, H. & Shibaev, V.P., Z. Phys. Chemie (Leipzig), 1983. 1. 161. 115. Kresse, H. & Shibaev, V.P., Makromol. Chern. Rapid Commun., 1984,5, 63. 116. Kresse, H., Tennstedt, E. & Zentel, R, Makromol. Chern. Rapid Commun.,

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118. Zentel, R., Strobl, G.R. & Ringsdorf, H., in Recent Advances in Liquid Crystal Polymers, ed. L.L., Chapoy, Elsevier, 1985, Ch. 17.

119. Tennstedt, E., Kresse, H. & Zen tel, R., Acta Polymerica, 1986,37, 685. 120. Zentel, R. & Wu, l, Makromol. Chem., 1986, 187, 1727. 121. Bormuth, F.1., Haase, W. & Zentel, R., Mol. Cryst. Liq. Cryst., 1987,148, 1. 122. Haase, W., Pranoto, H. & Bormuth, F.1., Ber. Bunsenges. Phys. Chem., 1985,

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1986,27, 185. 138. Attard, G.S. & Williams, G., Polymer Commun., 1986, 27, 2. 139. Attard, G.S. & Williams, G., Polymer Commun., 1986,27,66. 140. Attard, G.S. & Williams, G., J. Mol. Electron., 1986,2, 107. 141. Attard, G.S. & Williams, G., Liq. Cryst., 1986, 1, 253. 142. Attard, G.S., Mol. Phys., 1986,58, 1087. 143. Attard, G.S. & Araki, K., Mol. Cryst. Liq. Cryst., 1986, 141,69. 144. Araki, K. & Attard, G.S., Liq. Cryst., 1986, 1, 301. 145. Attard, G.S., Araki, K. & Williams, G., J. Mol. Electron., 1987,3, 1. 146. Attard, G.S., Moura-Ramos, lJ. & Williams, G., J. Polymer Sci.: Part B:

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Chem., 1989, 190, 2463.


152. Kresse, H., Wiegeleben, A. & Kriicke, B., Acta Polymerica, 1988, 39, 583. 153. Simon, G.P., Polymer, 1989, 30, 2227. 154. Massalska-Arodz, M. & Kresse, H., Liq. Cryst., 1990, 7, 361. 155. Attard, G.S., Williams, G. & Eawcett, A.H., Polymer, 1990,31,928. 156. Araki, K., Polymer J. 1990, 22, 546. 157. Kapitza, H., Zen tel, R., Twieg, R.J., Nguye, e., Vallerien, S.u., Kremer, F. &

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1987, 188, 1501. 159. Kremer, F., Vallerien, S.u., Zentel, R. & Kapitza, H., Macromolecules, 1989,

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Cryst., 1990, 7, 217. 161. Vallerien, S.u., Zentel, R., Kremer, F., Kapitza, H. & Fischer, E.W.,

Makromol. Chem. Rapid Commun., 1989, 10, 333. 162. Vallerien, S.u., Kremer, F., Fischer, E.W., Kapitza, H., Zen tel, R. & Poths,

H., M akromol. Chem. Rapid Commun., in press. 163. Vallerien, S.u., Kremer, F., Kapitza, H., Zentel, R., Scherowsky, G., Fischer,

E.W., Proceedings of the 13th International Liquid Crystal Conference, Vancouver, Canada, 1990.

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165. Heinrich, W. & Stoll, B., Colloid & Polymer Sci., 1985, 263, 873. 166. Adam, G. & Gibbs, 1.H., J. Chem. Phys., 1965,43, 139. 167. Kapitza, H. & Zentel, R., Macromol. Chem., 1988, 189, 1792. Suzuki, T.,

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168. Bar, Ch. & Heppke, G., Mol. Cryst. Liq. Cryst. Lett., 1986,4, 31. Yoshino, K., Ozaki, M, Kishio, Sh., Sakurai, T., Mikami, N. & Honma, M., Mol. Cryst. Liq. Cryst., 1987, 144, 87. Mohr, K., Kohler, S., Worm, K., Peltz, G., Diele, S., Zaschke, H., Demus, D., Andersson, G., Dahl,!., Lagerwall, S.T., Skamp, K. & Stebler, B., Mol. Cryst. Liq. Cryst., 1987, 146, 151.

Chapter 5

Lyotropic Side Chain Polymer Liquid Crystals

P.J. Hall and G.J.T. Tiddy Unilever Research, Port Sunlight Laboratory, Quarry Road East,

Bebington, Wirral, UK, L63 3J W

5.1 INTRODUCTION

The incorporation of liquid-crystal-forming molecules into large polymeric structures has opened up a wide and diverse new field in both organic and physical chemistry, This new interest has centred mainly on the synthesis, physical properties and technical applications of thermotropic polymer liquid crystals (PLCs), about which much will be mentioned in this book in Chapters 3 and 7. Lyotropic PLCs too have found many useful applications, for example in the production of new materials such as 'Kevlar' and other polyamides. 1 ,2 However, a further subclass of these polymeric materials, side chain lyotropic PLCs, has received comparatively little attention. The majority of the examples of such structures contain surfactant molecules as the monomeric side chains. These surfactant moieties are composed of a hydrophobic and a hydrophilic part and as such may be termed 'amphiphilic'. In the monomeric state surfactants may aggregate into micellar structures which, in concentrated systems, may further pack together to form one or more liquid crystal structures. This phenomenon highlights a major difference between thermotropic and lyotropic side chain liquid crystals: in the former it is the packing of molecules which gives the degree of ordering required, whilst for the latter it is the packing of the micellar aggregates which gives rise to liquid crystal formation. Polymerization may affect this process in a number of

237

238 P.J. Hall and G.J. T. Tiddy

ways, for example by altering the size and shape of micelles 3 or even by preventing micellization altogether. Here we might predict that the degree of polymerization of the polymer chain may be influential, since for the majority of micellar aggregates there is an upper limit to the aggregation number. Furthermore, variation in the polymer type and points of attachment of the side chains and head groups makes possible a large variety of PLC structures. To date only a few have been studied.

In the initial studies of compounds likely to form lyotropic side chain PLCs, the main interest was focused on their micellization properties in dilute solution. 4 ,5 An early study on the effects of polymerization in a liquid crystal phase was reported by Friberg et al., 3 but it was not until the work of Finkelmann et al. 6 that the effect of polymerization on liquid crystal formation was studied systematically. Analogous work on thermotropic systems 7.8 had indicated that polymerization brings about an extension of the mesophase region over a wide range of temperature. Finkelmann produced evidence suggesting that in lyotropic systems the same effect was occurring. In addition, polymerization appeared to extend the meso phase region over a broader composition range, this being an additional feature of two-component lyotropic systems.

More recent work 9 ,lo has shown that the distinction between lyotropic and thermotropic polymer liquid crystals need not be so rigidly defined. If the amphiphilic side chains have an element of rigidity built into them, using for example a biphenyl group, then some of the meso phases formed closely resemble those of thermotropic side chain polymers. Furthermore, polymers which produce lyotropic liquid crystals may well form mesophases in the absence of a solvent, should the molecular structure favour them. (Note that many surfactants, particularly ionic amphiphiles, also form thermotropic mesophases.) Some specific examples will be discussed in due course.

In this chapter we will first review the principles of liquid crystal (LC) formation for small-molecule amphiphiles. This will include a few of the techniques used to characterize such LC phases. We will then describe the synthetic routes to polymeric surfactants, and the likely effects of structural constraints on mesophase architecture. Most frequently the polymerizable group is a carbon~carbon double bond which may be induced to react with its nearest neighbour by thermal or photochemical initiation. This can take place when the molecules are in a micellar or liquid crystal state or even at an air~water interface. The properties of the resultant polymer will obviously be influenced by the choice of technique.

Lyotropic Side Chain Polymer Liquid Crystals 239

When polymerization is effected at the air~water interface, monolayer polymer films may result. Such films have attracted a great deal of attention in recent years by virtue of their similarity to biological membranes.!! Polymeric films do not really fall within the strict definition of liquid crystal polymers. However, the molecular arrangement of the side chains resembles that found in many liquid crystal phases, most notably the lamellar phase, and could therefore realistically be described as one half of a lamellar sheet. Bilayer films, on the other hand, may be thought of as a single lamellar layer, and therefore the comparison becomes even more appropriate. The technical applications of polymer membranes are extremely diverse and for these reasons a short section will be devoted to their history, their uses and their current status.

It is very often the commercial interest in novel materials which stimulates the growth in their study and eventual exploitation. This is certainly true in the case of thermotropic liquid crystals and their application in electro-optic displays.!2 Indeed, the production of highstrength, high-modulus fibres! has seen a wealth of interest in lyotropic main chain polymers. The use of lyotropic side chain polymers has, by comparison, been less well publicized. This is not to say that there are no applications. Alkyl polyoxyethylene surfactants attached to polysiloxane polymers have found uses in many personal care products such as liquid soaps, shampoos, skin creams, and hair mousses.13 Unfortunately the physical properties of these and other similar materials have been closely guarded secrets and the amount of information available in the literature is low. The limited data which does exist, however, provides us with some interesting structure ~ behaviour relationships.

5.2 PHYSICAL PROPERTIES OF SURFACTANTS

5.2.1 Dilute Micellar Solutions Chemical compounds known as surface active agents, or surfactants, consist basically of two parts: a hydrophilic 'head group' and a hydrophobic 'tail', as illustrated schematically in Fig. 5.1. The hydrophobic portion, which may be linear or branched, interacts only very weakly with water molecules. Moreover, the strong attractive interactions between the water molecules arising from dispersion forces and hydrogen bonding act cooperatively to squcczc thc hydrocarbon out of the water (the hydrophobic effect). The head group, which consists of a strong bipolar or ionic portion, interacts with water via dipole~dipole or

240 P.J. Hall and C.J.T. Tiddy

hydrophilic , head-group'

hydrophobic , tail'

Fig. 5.1. Schematic illustration of a surfactant molecule.

ion-dipole interactions and is solvated. Consequently this part of the molecule is said to be hydrophilic. Owing to the different environmental preferences of these two distinct regions, such molecules are also termed 'amphiphilic'. Their classification puts surfactants into various groups depending on the nature of the head group, that is, anionic, cationic, non-ionic or zwitterionic. Typical examples of these compounds are as follows:

Anionic: Cationic:

Non-ionic:

Z witter ionic:

Sodium dodecyl sulphate (SDS), C 12 H 2SSOiNa+ Dodecyl trimethylammonium bromide (DT AB),

C12H2SN+(CH3hBr-Dodecyl hexaethylene glycol monoether,

ClzHzs(OCHzCHz)60H 3-Dimethyl dodecylamino propane sulphonate,

ClzHzsN+(CH3hCHzCHzCHzS03

The structure of the hydrophobic tail is usually limited to either hydrocarbon or fluorocarbon chains. Compounds possessing a single such tail are the most common, although molecules with several linear or branched tails are known. Dialkyl surfactants such as lecithin (below) are particularly important naturally occurring materials.

0-

I + CHz-O-P-O-CHz-CHz- N(CH3h

II o Lecithin

Most, if not all, surfactants are soluble in water, but compounds with an alkyl chain length of C 16 or greater become increasingly insoluble. The physicochemical properties of aqueous surfactant systems are a consequence of the tendency of the non-polar groups to avoid contact with


water at the same time as the polar part tends to be strongly hydrated. The adsorption of surfactants at interfaces is one such example of this and is particularly important in the context of polymerization of adsorbed layers at the air-water interface, discussed more fully at the end of the chapter. Another example is the extensive aggregation of amphiphilic molecules into structures known as 'micelles'. The concentration at which this process occurs is known as the critical micelle concentration, usually abbreviated to CMC. The driving force behind micellization is again the hydrophobic effect 14 which dictates that contact between water and hydrocarbon be kept to a minimum. Micelle formation enables the polar head groups, which confer solubility on the monomeric surfactant, to remain hydrated.

Each surfactant has a characteristic CMC value at a given temperature, which may be determined by monitoring the change of a particular physical property as a function of surfactant concentration. The physical properties most commonly monitored include surface tension, conductance, vapour pressure and turbidity. Figure 5.2 shows the dependence of surface tension on concentration in the region of the CMC for the non-ionic surfactant dodecyl hexaethylene glycol monoether. Below the CMC increasing surfactant concentration leads to a steady decrease in surface tension (y). At the CMC micellization becomes energetically more favourable and further addition of

50

40

30

20

-6 -5 -4 -3

log (concentration I mol dm-3)

Fig. 5.2. y vs log c plot for n-dodecyl hexaethylene glycol monoether in water at 25°C.

242 P.l. Hall and G.l. T. Tiddy

surfactant results in an increase in the number of micelles. Thus the monomer concentration and consequently the surface tension remain fairly constant. As a rough working approximation it is useful to remember that many C12 ionic surfactants have CMC values of approximately 10- 2-10- 3 mol dm - 3. These values decrease with increasing alkyl chain length. The CMC values of non-ionic surfactants are usually lower than those of ionic materials of comparable chain length. For example, the CMCs for dodecyl polyethylene glycol monoethers (C 12EOn) are in the range of about 10- 4_10- 5 mol dm- 3. The larger value for ionic surfactants is due to unfavourable electrostatic repulsions between the charged head groups at the micelle surface. The CMC is extensively used as a very informative characterization of the self-association of surfactants. Mukerjee and Mysels have produced a valuable compilation of CMC data 15 which also includes a critical discussion of the various techniques available for their determination.

In 1936 G.S. Hartley suggested that the properties of solutions of ionic surfactants could be explained by the formation of one type of micelle. 16

This is shown schematically in Fig. 5.3. The micelle is spherical in nature with a radius approximately equal to the length of a hydrocarbon chain. The aggregation number is in the range 50-100 monomeric units. Hartley also considered that many of the counterions were bound to the head groups in order to alleviate the high charge density at the micelle surface. This factor explained a drop in conductance on association. The structure proposed by Hartley was correct from many points of view, though more recent work has indicated that rapid molecular motion causes the surface

o o

o o

Fig. 5.3. Schematic representation of a micelle formed by an ionic surfactant. Water (not shown) surrounds head-groups and counterions. Dotted lines mark

approximate boundary of polar and non-polar regions.


of the micelle to be slightly rougher than expected.! 7 Moreover, many modern techniques such as light scattering indicate that in many systems the micelles are not spherical but can be rod- or disc-shaped/ 8 as depicted in Fig. 5.4. It should also be mentioned that, although micelles are commonly presented with the alkyl chains pointing inward towards the centre of the micelle, the chains are in fact highly disordered and even the terminal methyl group may exist at the micelle surface for a finite period of time. On the other hand, the notion that water molecules might penetrate the micellar interior has been largely disproved in recent years.!9

(] rm Wlj i~J Ilw IllJ} j~) (a)

(b)

Fig. 5.4. Schematic representations of (a) rod- and (b) disc-shaped micelles.

5.2.2 Micelle Size and Shape Information on micelle size and shape can be obtained from a variety of techniques such as light scattering, neutron scattering, viscosity measurements and osmotic pressures. Typically the micelles have a closely spherical shape in a rather wide concentration range above the CMC. Often there is no great change in shape until the surfactant solubility limit is reached, when a liquid~crystalline phase normally separates out.

A qualitative description of the factors responsible for micelle shape has been put forward by a number of different workers.!4.2o,2! At the micelle surface there exists a balance of forces between head group repulsions and hydrocarbon~water repulsions. The former are particularly prevalent for surfactants containing ionic head groups. Head group repulsions serve to increase the average surface area per molecule (a). The


hydrocarbon chain-water repulsions tend to reduce a. An additional contribution comes from the different energies corresponding to either trans or gauche conformations in the bonds of the alkyl chains. The overall equilibrium at the micelle surface may be described by eqn (6.1):

o C(n) JiN=ya+-

a (5.1)

where Ji~ is the free energy per surfactant molecule in a micelle composed of N molecules, y is the surface tension at the hydrocarbon-water interface, and C(n)/a is a term which describes the head group repulsions.

In addition to the above there is a further constraint to be applied. This is based on the assumption that the micelle radius cannot be longer than the all-trans length of the hydrophobic chain (I). Thus there is a maximum volume for spherical and cylindrical micelles which in turn corresponds to a minimum surface area per molecule. Assuming a smooth micelle surface the two values for the different shapes are given by:

Sphere

Cylinder

a(min) = 3v/1

a(min) = 2v/1

where v = volume of hydrophobic chain. For a C 12 hydrocarbon chain, with a maximum length of 15 A and assuming the volumes of CH2 and CH3 groups to be 27 A 3 and 54 A 3, respectively, this gives

a(min)=70A2 (sphere) a(min)=47 A2 (cylinder)

The alkyl chain conformational freedom (gauche/trans isomerization) may also have an effect on the limiting values of a. The all-trans conformation becomes less favourable for long chains, which effectively reduces the surfactant chain length and increases the limiting values of a. Finally, there may be a contribution from the thermal fluctuations (kT) in the system which can bring about variations in micelle size and shape. Because of these fluctuations the micelle surface is not completely smooth, hence the maximum micelle radius may be slightly larger than that suggested by the 'packing constraints' model. The values of a(min) given above should be regarded as guides rather than exact limits.

It is important to know the micelle aggregation number because this influences the micelle shape, micelle kinetics, and transport properties. Translating the micelle size into the number of monomers in the micelle requires various assumptions the validity of which is difficult to test


experimentally. The aggregation number of a micelle can be obtained in a variety of ways. Firstly, the number of micelles may be 'counted' using a spectroscopic technique and, with the additional information about the amount of surfactant in micelles, the micelle aggregation number may be obtained; this forms the basis of the fluorescence quenching method which, for example, gives about 70 monomers in a sodium dodecyl sulphate (SDS) micelle. 22 Secondly, a spectroscopic or other property of the surfactant which is distinctly different in the monomeric and micellar states may be monitored as a function of concentration. There are several possibilities but the 13C NMR chemical shift is probably the most generally and conveniently studied parameter.23 From the directly determined aggregation numbers, the radius of a minimal sphere may be calculated from the considerations of micelle volumes which we have just considered. In general there is a close agreement between the values obtained via this method and those arrived at by direct measurement. This implies a very distinct separation of alkyl chains and water, with no water penetration into the micelles.

Micellar growth with increasing concentration is most obviously demonstrated by a concomitant increase in viscosity. It is now well established that micellar growth leads to long rod-like aggregates and is favoured by decreasing temperature, addition of electrolyte and increasing surfactant chain length. The nature of the counterion in ionic systems is also an important consideration, with strong counterion binding favouring long rod micelles.

5.2.3 liquid Crystal Formation of Small-Molecule Surfactants

As the surfactant concentration is increased above the CMC the number of micelles is also increased, and this ultimately leads to an increasing extent of repulsive interactions arising from excluded volume, hydration, steric and electrostatic effects. For micelles of ionic surfactants the strongest interaction is the electrostatic force. This acts at a distance from the micelle surface which varies with the electrolyte concentration. For non-ionic surfactants the interactions are much weaker and are limited to repulsions arising from the size and solvation of head groups. Steric effects extend well away from the alkyl chain-water interface for surfactants with large head groups such as alkyl polyoxyethylene compounds. The fact that each polyoxyethylene chain is associated with a certain amount of water increases the range. This is the so-called hydration force. The increased interactions between micelles as the concentration of

246 P.J. Hall and G.J.T. Tiddy

surfactant increases has several effects. Firstly, the micelles tend to grow in size. The value of a decreases, sometimes resulting in a sphere/rod or a rod/disc shape alteration. As the micelles become bigger and more numerous, a transition from a disordered state to an ordered array is produced, the micellar solution becoming a lyotropic liquid-crystalline phase.

In two-component surfactant-water systems there are three major types of liquid crystal structure. Their structures have been elucidated over the years using primarily low-angle X-ray diffraction,24 but with other evidence from optical and electron microscopy and NMR spectroscopy.25 The mesophase types have been labelled according to various conventions but the system adopted by us previously25 will be used here. The first type we shall consider is the lamellar (L,) phase. It consists of ordered bilayer micelles separated by water layers which may vary in thickness from 10 to 100 A depending on the nature of the surfactant. This is shown schematically in Fig. 5.5. The micelles possess an essentially fluid interior resembling that found in liquid hydrocarbons of corresponding chain length. However, this liquid character is restricted by the requirement for polar head groups to reside at the alkyl chain-water interface.

The second type of meso phase encountered is the hexagonal phase, labelled HI or Hz. The subscripts denote normal hexagonal and reversed hexagonal phases, respectively. Reversed (oil-continuous) phases are formed by surfactants with large hydrophobic chains and small head groups in non-polar solvents or very concentrated aqueous solutions. The aggregation process is driven by attractions between polar groups and water which make up the micelle interior. The normal hexagonal phase (water-continuous), HI> is depicted in Fig. 5.6. It is made up of hexagonally packed rod-like micellesz4 whose surface separations are typically in the range 10-50 A. For a surfactant which displays both normal hexagonal and lamellar mesophases, the hexagonal phase is always found at lower surfactant concentration. This can be predicted on theoretical grounds in that the rod-like aggregates expose more surface area to the surrounding water (i.e. a larger a value).

The third type of liquid crystal structure is the cubic phase, of which there are two distinct classes. Unlike the well-characterized structures of the lamellar and hexagonal phases, the cubic structures have yet to be fully elucidated. The cubic phases are labelled II, Iz and Vb Vz, the subscripts I and 2 again corresponding to normal and reversed structures. Normal (I tl and reversed (Iz) cubic phases simply consist of small, often spherical micelles arranged in various arrays. One very common


water

head-groups

Fig. 5.5. Schematic representation of lamellar phase (L,).

structure consists of short rod micelles, but still arranged in a cubic array.26 Figure 5.7 shows a representation of a body-centred cubic structure. These phases, when they occur, form at compositions between isotropic micellar solution, labelled Ll and L2, and hexagonal phase (HdH2)' The second group of cubic phases (V dV 2) is formed at compositions between lamellar and hexagonal phases.2729 The structure is thought to be composed of interconnecting three-dimensional networks of surfactant molecules. A recently published representation of this phase28 is shown in Fig. 5.8. The centres of the four circular openings of the unit (Fig. 5.8(a)) may be imagined as forming the vertices of an inscribed tetrahedron. The walls of the unit consist of a bilayer (inset) with water on either side of the wall. Several different structures are possible, though further work is necessary to fully elucidate the possibilities. The surface area per molecule (a) in this structure should fall somewhere between that of lamellar and hexagonal phases since a

248 P.1. Hall and G.J. T. Tiddy

alkyl chains

Fig. 5.6. Schematic representation of normal hexagonal phase (H 1)'

Fig. 5.7. Schematic representation of II body-centred cubic structure.

decrease in a leads to the sequence: cubic (I d --> hexagonal (H d --> cubic (V 1) --> lamellar (L,l --> reversed phases.

All the common meso phases are viscous, with viscosity usually increasing in the order:

The range covered is enormous. Lamellar phases often flow under gravity, while the 11 phase can be so rigid that it is difficult to insert a spatula into the material. A wide variety of other meso phases can occasionally occur. Some have yet to be fully identified and their


(a) (b)

Fig. 5.8. (a) Schematic representation of the repeating unit of a bicontinuous cubic phase (V d. (b) Stacking of repeat units. (From Ref. 28 with permission of

the American Chemical Society.)

structures have so far only been postulated. One such mesophase that is worthy of note is the lyotropic nematic phase, which has received intensive study only relatively recently.30.31 It occurs in a minority of systems between Ll and HI or betwen Ll and La for molecules in which the hydrocarbon chain is usually fairly short (n < 14). The nematic phase may be uniaxial or sometimes biaxial. Structurally the former is composed either of ordered small cylindrical micelles (Nc) or ordered small disc-shaped micelles (ND). Its viscosity is the lowest of all lyotropic mesophases.

The composition and temperature ranges of the various mesophases in surfactant~water systems may vary quite considerably. The relationship is most conveniently expressed using phase diagrams. These depict the areas and boundaries of phase stability in the form of a graph, where the axes represent compositions and/or temperature. A typical example of a binary phase diagram is shown for the dodecyl hexaethylene glycol monoether (C12E06)~water system in Fig. 5.9. This surfactant displays large regions of both lamellar and hexagonal mesophases together with a narrow band of cubic V 1 phase located between them. 3 2 Also shown is the lower consolute boundary or 'cloud curve'. The cloud curve represents the boundary between isotropic micellar solution (Ld and a two-phase region (Ll + W) corresponding to phase-separated concentrated and dilute surfactant solutions. This phase separation is a consequence of decreasing surfactant solubility with increasing temperature, probably brought about by conformational changes in the polyoxyethylene head group which disfavour hydration.


90

W . L,

60

0" f:

30 L 1

5

o 25 50 75 100 composition (wt (Jr. (' 12 to,)

Fig. 5.9. Phase diagram of the C12 E0 6-water system.

5.2.4 Techniques of Characterization The structures of the various lyotropic mesophases mentioned so far have been elucidated over the years primarily using low-angle X-ray diffraction. 24,33 An X-ray diffraction pattern of a liquid crystal provides information not only on the state of organization of the hydrocarbon chains but also on the crystallographic lattice of the micellar structure. It must be emphasized, however, that often the X-ray method alone cannot define the absolute structure of a liquid crystal phase because too few diffraction lines are observed. In these cases, a knowledge of the position and extent of the mesophase region in the phase diagram, measurements by other techniques (NMR, optical microscopy), and information such as the size, shape and chemical nature of the surfactant are necessary before a reliable identification can be made.

The diffraction pattern obtained for a lyotropic liquid crystal is characterized by a series of sharp reflections at low angles corresponding to interlayer spacings ranging from about 10 to 100 A and a wide-angle diffuse reflection corresponding to a spacing of 4·5 A. The number of reflections depends on the amphiphile, the meso phase type and the composition. Between 2 and 7 reflections are commonly encountered, although the intensity of the higher-order lines is usually low. The pattern of reflections varies according to the symmetry of the mesophase (1:1/2:1/3..., lamellar; 1:1/.J3:1/.J4 ... , hexagonal, etc.).24,29


Where the structure of the mesophase can be positively identified, the positions of the observed reflections, together with other information such as composition, density and geometry of the phase, can be used to calculate the dimensions of the structure and in particular the surface area per surfactant molecule (a) at the alkyl chain- water interface. Perhaps the quickest way to characterize a lyotropic mesophase is by polarizing optical microscopy. Lamellar, hexagonal and nematic phases are optically anisotropic and all exhibit a variety of characteristic textures. The lamellar phase textures are generally of the focal-conic type, the units of which often appear as tiny 'Maltese Crosses' or 'spherulites'. These may be encountered individually or as the units which make up 'mosaic' textures (Fig. 5.10). The textures most commonly observed in hexagonal meso phases are focal-conic 'fan-like' textures (Fig. 5.11) or rather featureless 'non-geometric' textures. These, together with the lamellar phase textures, have been reviewed extensively by Rosevear. 34 Nematic textures are typically analogous to those of thermotropic nematic systems i.e. Schlieren texture.

With regard to the cubic phases that are optically isotropic, identification by optical microscopy can be made owing to the very

Fig. 5.10. Typical 'mosaic' texture of lamellar phase containing spherulites, sodium octanoate, 50% water, at room temperature.

252 P.J. Hall and G.). T. Tiddy

Fig. 5.11. Typical 'fan-like' texture of hexagonal phase, CI2 EOs, 40% water, at room temperature.

high viscosity. A distinct band between HI and L. meso phases in a concentration gradient of the surfactant-water system is strong evidence for a V 1 phase. Similarly, the occurrence of such a band between L 1 and HI is good evidence for the existence of an II phase.

A qualitative picture of the range of meso phases occurring for a particular surfactant may be obtained by conducting a penetration experiment. 35 This involves placing a small amount of surfactant between a glass slide and cover-slip followed by application of one or two drops of water to the edge of the cover-slip. The water penetrates between the slide and cover-slip, establishing a concentration gradient at the surfactant-water boundary. This may then be viewed on the microscope as a function of temperature. The onset or disappearance of the different phase regions can then be monitored. Although not quantitative, penetration experiments provide a rapid assessment of surfactant phase behaviour which may be carried out using very small amounts of material. A second technique (preferably X-ray diffraction) should always be employed to confirm the meso phase structures.


5.3 SYNTHESIS OF LYOTROPIC SIDE CHAIN POLYMER LIQUID CRYSTALS

As we have already seen in Section 5.2, monomeric amphiphilic molecules contain two distinctive regions. Incorporation of such molecules into side-chain polymers may in principle be brought about by polymerization at the hydrophobic or hydrophilic ends, thereby giving rise to the structures shown in Fig. 5.12. In the type A polymers the pendant side chains extend from a hydrophobic backbone, whereas type B polymers are formed by amphiphiles with polymerizable hydrophilic groups. The micellar aggregates of these two different types may be expected to be quite different, as has been suggested by pioneers in the field. 3 •36 More

Type A Type B

Fig. 5.12. Schematic representation of two categories of polymeric amphiphile.

recently the synthesis and properties of novel star-block liquid crystalline copolymers have been described 37 in which the lyotropic side chains radiate from a central hydrophobic moiety, giving rise to disc-like molecules. The idealized structure of such polymers (Fig. 5.13) bears some resemblance to the discoidal amphiphiles synthesized by Boden et al. 38

and by other groups of workers. 39 .4o The central hydrophobic 'discs' are typically triphenylene, benzene, copper phthalocyanine or cyclotriveratrylene. The R groups which project from the core may vary greatly in type and in number, though molecules based on benzene or triphenylene contain a maximum of six. Whether or not these types of compounds could be described as polymers, arguably they are at least oligomers. They appear to be more commonly referred to, however, as multipolar amphiphiles,40 which is perhaps a more acceptable term. Nevertheless, these compounds are structurally very interesting and in the present context are worthy of inclusion.

Returning to the structures in Fig. 5.10, polymerization may be carried out most effectively by one of two methods. Both these approaches rely on the polymerization of terminal vinyl groups at the end of the


Lyotropi c block

Flexible block

Fig. 5.13. Topology of lyotropic star-block polymer liquid crystals.

hydrophobic chains. In the simplest case this reaction proceeds via a free-radical process which may be initiated thermally or by using a source of y-radiation. Larrabee and Sprague41 showed that for the polymerization of dilute aqueous sodium lO-undecenoate the initial concentration of monomer must exceed the CMC otherwise reaction cannot occur. The explanation for this behaviour is simply that in the micellar environment the reactive species are localized. Rate of reaction is conveniently monitored by disappearance of NMR absorption due to the protons adjacent to the carbon-carbon double bond. Friberg et al.42 found that polymerization of an HI phase of sodium undecenoate-water at 60°C yielded a polymer composed of approximately 270 amphiphilic chains. Thus under conditions of closely packed alkyl chains quite high degrees of polymerization (within the limit of the aggregation number) may be obtained.

Other workers have used initiators to assist in the polymerization.43

Such initiators may be obtained thermally (e.g. azobisisobutyronitrile, AIBN), or photochemically (e.g. 2,2-dimethoxy-2-phenylacetophenone, DMPAP). Though only sparingly soluble in aqueous solution, the initiators preferentially reside in the micelles and therefore need only to be added in relatively small amounts. In reporting the polymerization of novel quaternary ammonium (meth)acrylates, Hamid and Sherrington43

have cast doubt on the possibility of producing a topochemical polymerized micelle because of rapid monomer exchange between micelles. Moreover, they have suggested the presence of oligomeric species which display micellar-like physical properties.

Polymerization by any of the above processes inherently yields a polymer having a fairly inflexible poly (meth)acrylate backbone. As we


shall see, such a structure confers significant translational and rotational restriction on the organizational capabilities of the resulting compound. An alternative method of producing comb polymers makes use of polysiloxanes, for the main chain component, which have much greater flexibility. These will be encountered in Chapter 7, but their favourable orientational properties are also suitable for lyotropic systems. The synthesis of such polymers occurs via the smooth addition of vinyl monomers to poly(oxymethylsilylene):

[ CH3 1 [ CH

3 1 ···-fi-O-... +CH2=CH-R4H-fi-O- ...

H n CH2-CH2-R n

R = amphiphilic group

The first example of this addition reaction using amphiphilic monomers was described by Finkelmann et a/. 6 in 1982. Their method, in common with most workers in similar fields, employed hexachloroplatinic acid as the catalyst.44 The progress of such reactions is usually followed by monitoring the decrease in intensity of the Si-H infrared adsorption45 at 2140 cm -1. Alternatively, the reaction may be monitored more quantitatively using 1H NMR spectroscopy.46 Similar chemical shifts of the Si-H and CH 2=CH- groups, however, can complicate an accurate interpretation. The olefinic compound is usually added in a 10% molar excess to ensure full reaction of the Si - H sites. This presents the problem of isolating the polymer in a pure state after the reaction. Provided purification can be carried out successfully, then fairly reproducible results should be obtainable. Detailed studies of the hydrosilylation reaction together with its shortcomings have been given by Gray et al.47 and also by He et al.48 In the latter paper the possibility of a competitive side reaction resulting in cross-linking at the terminal dimethylsilyl functions is considered. This does not take place, however, when the side chain mesogen is added in excess.

Only one example of the type B polymers shown in Fig. 5.12 has appeared in the literature.49 This material, synthesized by Jahns and Finkelmann, was produced using radical polymerization of an amphiphile containing a vinyl moiety at the terminal end of the hydrophilic

256 P.l. Hall and C.l.T. Tiddy

head group. The phase behaviour of this material together with further examples of polymeric amphiphiles is discussed in the following section.

5.4 PHASE BEHAVIOUR OF LYOTROPIC SIDE CHAIN POLYMER LIQUID CRYSTALS

Only relatively recently have lyotropic polymers been investigated with respect to their liquid crystal properties. In 1982 Finkelmann, Luhmann and Rehage synthesized a series of surfactants having lO-undecenoic acid as the hydrophobic region and a polyethylene glycol moiety as the hydrophilic region.6 The corresponding polymers were produced by addition of these surfactants to a polysiloxane main chain. The structure of the polymers is shown below:

This represents a major advance over previous work because polymers of well-defined molar mass could be obtained and because the hydrophobic and hydrophilic moieties are well separated.

The phase behaviour of both the monomer and polymer having m = 8 was determined by polarizing optical microscopy and the resulting phase diagrams were found to differ quite considerably. The issue of whether the polymers would display any liquid crystal behaviour at all was quickly resolved. In fact, attachment of mesogenic side chains to silicon sites of a polysiloxane backbone does not restrict the conformational freedom of the amphiphiles beyond that necessary for the packing of chains into LC arrays. Moreover, at first glance it would appear that quite the reverse occurs. Fig 5.14 shows the phase diagram of the monomeric surfactant-water system,6 where the monomeric surfactant has the structure CH2=CH(CH2)SCOO(CH2CH20)sCH3' This surfactant displays a single hexagonal liquid crystal phase in a concentration range from 49% to 70% surfactant. At 53% monomer concentration the phase reaches an upper temperature limit of 19°C. At this concentration three water molecules per oxygen atom on the glycol chain are available.


00

eo o

d p

cD ,,-0 0 .---0- __ 0

()-¢- -<>- -

LO

20

-20

20 40 60 eo

MONOMER CONTENf %

Fig. 5.14. Phase diagram of the monomeric surfactant- water system.

In contrast, the phase diagram of the polymeric surfactant shown in Fig. 5.15 displays three well-defined liquid crystal regions, namely lamellar (L.), hexagonal (H I) and cubic (V I)' The two former phases are optically anisotropic and may therefore be easily identified by their characteristic textures. The narrow region of cubic phase, however, is optically isotropic and was identified by virtue of its high viscosity and shear birefringence.

It has been suggested that the enlarged meso phase regions displayed by the polymer- water system are a consequence of a stabilization effect brought about by restriction of motion of the polymer side chains. However, in the process of polymerization the alkyl chain of the surfactant is effectively lengthened by one monomer unit of the original polymer


'on

RO

50

LO

20

- --- -----~ -U-<>O-~

- 20 r

L----'2'r-"O --- ' -0 - ---,6'0- ---:<8'0:--~ POLYMfR C:) NTENT 'f,

Fig. 5.15. Phase diagram of the polymeric surfactant-water system. Black region represents bicontinuous cubic phase, V \.

(equivalent to about three CH 2 groups). Mitchell and co-workers have investigated the effect of varying alkyl chain length in alkyl polyoxyethylene compounds whilst maintaining the hydrophilic chain length constant. 32 Such studies indicate that increase in alkyl chain length leads to a broadening of the lamellar phase region, with little change in the hexagonal phase region. Accompanying these trends is a general shift of the cloud curve to lower temperature. With very short-chain surfactants (C n , n "" 8) meso phases do not occur at all, or form over very limited temperature and composition ranges. The presence of terminal vinyl moieties in surfactants is known to increase the CMC by an amount equivalent to shortening


the alkyl chain by about 1·1 CH2 units. 14 The smaller volume compared to a -CH2CH 3 moiety would also be expected to reduce mesophase formation. Hence the difference in meso phase formation between the monomer (Fig. 5.14) and polymer (Fig. 5.15) is that expected between conventional C g and C1CC 16 surfactants. This is only slightly more than the changes in hydrophobic volumes between monomer and polymer. There is no convincing case that polymerization per se has a large influence.

It is interesting to note that a V 1 phase is observed by Finkelmann. Recent work so has demonstrated that fairly minor conformational restrictions of an alkyl chain disfavour the formation of VI phases. Hence the VI phase would not be expected for polymeric surfactants. Indeed, in recent work, on an almost identical polymer with a degree of polymerization of about 50, we have been unable to identify any cubic VI phase, although large areas of HI and La were presentS 1. This leads us to tentatively suggest the presence of some monomeric surfactant in the first polymeric surfactants.

A novel variation on the type A polymer contains a rigid biphenyl moiety introduced into the hydrophobic part of the molecule.9 A series of such compounds in both monomeric and polymeric forms were recently studied. An example of the type of compound described, together with its phase diagram, is shown in Fig. 5.16. In addition to lamellar and hexagonal phases this material also forms a lyotropic nematic Nc phase which occurs over a narrow range just to the low concentration side of HI' This phase is believed to consist of relatively short orientationally long-range-ordered cylindrical micelles and its presence is promoted largely by the rod-like biphenyl group. As part of the same study, Luhmann and Finkelmann investigated the effect of varying the degree of polymerization (d.p.) of the polysiloxane backbone. They found that as the d.p. (r) increased a number of trends become apparent. Firstly the stability of the II-phase decreased with increasing r. Conversely, the thermal stability of the L. phase was found to increase while that of the HI phase showed no major change. Since spherical micelles (from which the cubic 11 phase is composed) are limited to a maximum aggregation number, as r increases there should come a point where 11 phases are no longer favoured. This effect has been used to partly explain the decrease in 11 stability with increasing degree of polymerization. Rod-like and bilayer micelles, however, are not restricted to an upper aggregation limit since their growth is permitted in one and two dimensions, respectively. Thus one would expect there to be no influence on phase behaviour by packing constraints. The increased thermal stability of the La phase

260 P.1. Hall and G.J. T. Tiddy

H,C-S,-ICK;l,o<:}Q-0I-C ,- CH,-Olo-CH, / \','A TF~ 9

'00

80

u 60

'" ~ 40 '<i a: UJ Q. :>' ;! 20

o

ONF ISOTROPIC -ONE A I SOT~OPI C lIOUID

20 40 60 BO ' 00

PO V'IER CON lEN- ....

Fig. 5.16. Structure and phase diagram of a polymer containing a rod-like moiety with water.

shown by the higher-molecular-weight polymers has therefore been attributed to change in interaction forces between neighbouring hydrophilic micellar regions.

A comparison of monomer and polymer phase diagrams has also been conducted for monosaccharide amphiphiles based on an N-methylglucamide: 52

o OH OH

~O ~~OH CH3 OH OH

The phase diagram of the monomer-water system exhibits only an L. phase in contrast to normal N-methylglucamide surfactants (up to at least Cd, which also show HI and V I phases. Unfortunately there is as yet no published work on such materials but our observations of high-curvature LC phases suggests that the hydrophilic acryloyl moiety in the above


structure loops round into the alkyl chain-water interface. Since curvature effects are weak in such surfactants, this looping of the hydrophobic tail constrains LC formation to La phase only. Polymerization of the acrylate moieties yields a material for which this cannot happen. Once again the La phase is observed on contact with water, this time being characteristic of normal longer-chain glucamides. The most interesting feature of the polymer phase diagram is the presence of a closed miscibility gap. As a cautionary note, however, this type of behaviour may be promoted by the presence of small amounts of impurities.

Among the simplest and best-studied of surfactants are the metal salts of long-chain carboxylic acids, because of their use in soaps. In fact the phase behaviour of many examples is well known in the literature particularly as a result of the pioneering work of McBain and co-workers53 in the first half of this century. It is not surprising, therefore, that the incorporation of such molecules into polymeric structures has already been investigated, 'polysoaps' being one such example. 54 The incorporation of sodium undecenoate into a side chain polymer has recently been described by several groups of workers.4 ,42,55 Typically, aqueous solutions of sodium 10-undecenoate at concentrations above the CMC are irradiated with ,-rays so as to induce polymerization. Using this technique, polymers having a d.p. in the region of 10 have been produced. 55 Micellization of the monomeric species effectively concentrates the reactant molecules and is necessary to facilitate the polymerization process. Most studies have focused on the micellar properties of the poly(sodium undecenoate) using intramicellar probes such as the fluorophore pyrene or a spin-labelled species. 55 Paleos et al. 4 have suggested, as a result of such studies, that the CMC of the polymeric micelle is equal to zero, i.e. a single polymer molecule constitutes a micelle in its own right which will maintain its identity to infinite dilution. While only monomers will be present at very low concentrations, some aggregation, albeit with low aggregation numbers (n ~ ca. 20), would be expected at higher concentration, for example adjacent to hexagonal phases or where long rod micelles could occur.

Although such micelles, at higher concentration, would probably pack together to form liquid crystal structures, no mention of any mesophase behaviour was made. An earlier study carried out by Thundathil et al.42

actually induced polymerization in the liquid crystal phase itself. Sodium undecenoate and water give a hexagonal phase in a concentration range of 44-55 wt% soap at 60°C. When ammonium persulphate (0'05 mol/litre) is present as the initiator, these concentration limits are shifted to 47% and 59% soap. Polymerization was carried out at 60°C and was monitored


using low-angle X-ray diffraction and polarizing optical microscopy. After 5 hours certain areas of the material had become isotropic. A further 36 hours' heating brought the reaction to completion. On cooling of the sample to room temperature a lamellar phase gradually appeared. Obviously the polymerization had caused the phase change. Analysis of the product polymer gave a relative molecular mass of around 52000, corresponding to a degree of polymerization in the region of 270. The change in phase behaviour was concluded to have arisen from volume contraction or the restriction of free rotation and translation of the end groups. However, a number of calculations involving packing constraint theory led the authors to predict only a 5% volume change and therefore to conclude that mobility restriction was the prime cause of the transition from cylindrical micelles to bilayers. Evidence from molecular modelling was provided to support this belief.

It is interesting to note that the above study considers polymers having a conformationally restricted polyethylene backbone, which obviously has a profound influence over the orientational capabilities of the molecules. Replacement of this rigid element by a flexible polysiloxane chain ought to relax the side chains in such a way so as to provide a far less restrictive structure. In fact, we have attached sodium undecenoate to a polysiloxane chain of approximately 50 repeat units. 56 The results have shown only a slight modification of phase behaviour, in accordance with a small chain lengthening of the hydrophobic chains.

Direct comparisons between hydrocarbon and siloxane main chains are few for lyotropic side chain polymer liquid crystals. Pietschmann et al. 57 have recently synthesised polymers of 1,3-diols which differ in the nature of the polymer backbone. DSC and optical polarizing microscopy were used to construct the phase diagrams. These were compared with the phase diagrams obtained for the analogous low-molar-mass compounds. The structures of the polyacrylate (I) and polymethacrylate (II) based materials are shown below:

R=H:(I) R=CH 3: (II)


The low-molar-mass compounds which are terminated by an olefinic group do not show the presence of any liquid crystal phase in the presence of water. If the olefin and ester groups are replaced by a simple hydrocarbon chain, however, then a lamellar phase is obtained over a fairly broad range of temperature and concentration. This implies that the increased hydrophilicity of the ester group is inhibiting liquid crystal formation. Again this material is behaving like an L/.-ill surfactant where both ends of the molecule can reside at the alkyl chain-water interface. As such it may well be behaving like a short-chain surfactant for which mesophase formation is inhibited.

For the polyacrylate (I) and the polymethacrylate (II) the lamellar phases exist over a greater temperature range. This is mainly due to the lower melting points induced by polymerization. The corresponding polysiloxane material is also reported to give a stable L. phase with or without water. Compared to the other polymers the crystal-L. melting point is slightly depressed while the L.-isotropic boundary is as much as 35°C greater. Thus it would appear that polymerization stabilizes the mesophase formed by the low-molar-mass diols. No phase changes are induced by the rigid polyacrylate or polymethacrylate, as was seen in the work of Thundathil et al.,42 though this is not surprising since conformational restrictions imposed by the hydrocarbon backbone favour the L. phase anyway. The extended phase boundaries observed for the polysiloxane were attributed to the greater flexibility of the siloxane main chain. However, the contribution to this effect made by effective alkyl chain lengthening was not considered.

In certain cases we have considered so far we have seen the effects of enforcing a particular shape or conformation on a micelle through the polymerization process. Therefore, the favoured aggregation of lowmolar-mass species may be greatly disturbed by the joining together of the alkyl chains. This effect is seen less for flexible polysiloxane chains than for the more rigid hydrocarbon-based chains. Similar conformational restriction and induction of LC phases has been demonstrated by Keller-Griffith et al.40 in the synthesis and characterization of multipolar amphiphiles attached to planar ring systems. These ring systems (benzene or triphenylene) constitute the hydrophobic cores of the molecules. The alkyl chains of the amphiphiles are radially attached to the rings by ester linkages with the hydrophilic cationic pyridinium or trimethylammonium head groups located at the opposite ends of the alkyl chains.

By analogy to thermotropic systems, 5 8 in which disc-like molecules stack themselves into columnar phases, it was hoped to show that the

264 P.J. Hall and G.J.T. Tiddy

synthesized compounds would form cylindrical micelles in the presence of water together with related liquid crystal phases at high concentration. A schematic representation of this concept is shown in Fig. 5.17. Indeed, in the majority of cases hexagonal phases built up by the close packing of the cylindrical aggregates were preferred over lamellar phases.4o For one series of compounds the packing of the disc-like molecules into cylinders was seen to be dependent on the number of polar head groups and the size of the central disc-like core. Another of the compounds appeared to form a lyotropic nematic phase to the low concentration side of the hexagonal phase. This phase is thought to consist of an anisotropic arrangement of finite anisotropic micelles. An interesting consequence of orienting the individual amphiphiles in a radial arrangement around a central core is that the ability to adsorb at planar interfaces becomes greatly restricted. Evidence for this is provided by plots of surface tension against concentration.40

* Disc-like

amphiphile Cylindrical micelle

Fig. 5.17. Schematic representation of the aggregation of disc-like amphiphiles into cylindrical micelles.

The surface tension of water is reduced only slightly by these molecules, reflecting the difficulty in orientation at the interface. Furthermore, the breakpoint in the y vs concentration plot corresponding to the CMC is not well defined and leads one to wonder what shape of 'micelles' could possibly be formed at the onset of micellization. This is a particularly good example of the interesting modifications in physical properties which may be induced by restricting the orientational freedom of amphiphilic molecules.

Also worthy of inclusion is the related work of Gallot and Douy/o who have successfully synthesized side chain polymers incorporating lipopeptides. Radical polymerization of terminal olefinic groups is once again the process used to produce the poly(meth)acrylamide


main chain. The monomeric precursors are shown by the general formula below:

R 0

I II HzC=C-C-NH-B-NH - (AA)p

where R=H or CH 3 , B=(CHz)n and AA=amino acid residue. The aqueous phase behaviour of the polymeric compound having n = 12 and p = 6·5 with polysarcosine as the hydrophilic head group was investigated. Two liquid crystal phases were observed, namely a lamellar phase (referred to as a smectic A (SA phase) and a nematic phase. Both phases were also observed in the dry material: at temperatures below 90"C the SA phase was present, while the nematic phase existed in the temperature range 90-150a C. Above this temperature the LC phase melted to an isotropic liquid. This polymer provides a good example of a material having both lyotropic and thermotropic properties and therefore goes a long way towards breaking the divisions between these two fields.

This work has recently been extended to investigate the effect of changing the amino acid residue. 59 .6o Early results show the expected link between head group size and meso phase structure. Hence polymers containing the bulky tyrosine side chain are able to form both La and HI phases, whereas alanine and glycine-containing polymers form only La meso phases. 60

Returning to the work of Finkelmann, a later paper describes the synthesis and phase behaviour of a surfactant having the polymerizable group located at the hydrophilic end of the molecule.49 Radical polymerization of this molecule

gives rise to a polymer of type B described in Section 5.2. Luhmann et al. 36 considered that polymers of this nature ought not to

be able to form liquid crystal phases of the normal type, i.e. watercontinuous phases. Instead they predicted that the constricted freedom of the head groups would lead to the formation of reversed phases. A comparison of the monomer and polymer phase diagrams obtained is shown in Fig. 5.18a and b. Both phase diagrams are complicated by the extremely broad miscibility gap which exists down to quite low temperatures. This is due to the reduced solubility of the polyoxyethylene chain

266

(a)

(b)

P.l. Hall and G.J. T. Tiddy

10 I'CI

80

TWO ISOTROPIC lIQUIOS

40

ISOTROPIC

ICE· CRYSTALLINE HYDRATE.

CRYSTALLINE HYDRATE CRYSTALLINE AMPHIPHllE

-20

20 40 60 60

MONOMER CONTENT ,n ",'Yo

100 l'CI

60

TWO ISOTROPIC LlaUIDS

60

ISO~

+-----~--.-------! ... T.OPI( SOLU

40 ISOTROPIC LIQUID" LAMELLAR liON

I L. ~ 20 f----......----il---,/\ _

ISOTROPIC LIQUID" HEXAGONAL ~

11H.·1"H,oro, E2

ISOTROPIC lIQUIO • CUSIC

ISOTROPIC LlQUIO" CRYSTALLINE HYDRATE

E, ~ ICE ·CRYSTALLINE HyDRATE ~

-20 i

100

20 40 60 60 100 POLYMER CONTENT In ... %

Fig. 5.18. (a) Phase diagram of the monomeric acrylate-water system. (b) Phase diagram of the polymerized acrylate-water system.


caused by the pr~sence of either the terminal acrylate group or indeed the polyacrylate cha:ln. In normal alkyl polyoxyethylene surfactants the head group owes its solubility to a relatively unrestricted freedom of conformation. Certain conformations are known to favour hydration by water more than others,61 and thus the restriction imposed by the rigid polymer chain may be contributing significantly to the lower solubility. Substitution of the terminal OH group by the polymerizable group also drastically reduces the hydrophilicity of such surfactants, as is known from work performed on methylated polyoxyethylene compounds. 62 Despite the broad miscibility gap present in both phase diagrams, liquid crystal regions are still encountered and these were ascertained to be of the normal and not reversed type. While the monomer appears only to form a hexagonal phase, the polymer forms in addition a cubic (V 1) and a lamellar phase, all of which show coexistence regions with the isotropic liquid. It is therefore concluded that suppression of normal liquid crystal aggregates which contain a high degree of surface curvature does not occur.

5.5 POLYMERIZATION IN ORIENTED MONOLAYERS AND VESICLES

Interest in this area appears to have developed at round the same time as for lyotropic side chain polymers. Certainly the types of molecules incorporated into the various polymeric structures bear a close functional resemblance to one another. The driving force behind studies in this field, however, has been the biological interest in producing systems which are able to mimic the properties of cell membranes. For a more in-depth introduction to the subject the reader is directed to other reviews, such as those by Ringsdorf et al. 6 3 and by Fendler. 64

Biological membranes typically contain about 50% each of lipids and proteins. Perhaps the most familiar model of the membrane make-up has been that of Singer and Nicolson65 (shown in Fig. 5.19). In it the lipids form a bilayer similar to those found in lamellar liquid crystals. The bilayer structure is spanned by large globular protein complexes which serve as selective barriers to the transportation of metabolites in and out of the cell. In addition to this they also confer a degree of stability on the structure, preventing it from breaking apart. The exact nature of this stabilization is obviously rather complex and extremely difficult to mimic in the synthesis of artificial membranes. Therefore, polymerization of the

268 P.l. Hall and G.J. T. Tiddy

Fig. 5.19. Schematic representation of the fluid mosaic model of a cell membrane.

lipid components has been studied as an alternative method for inducing long-term stability.

The polymerizable groups and methods of polymerization are exactly the same as described in Section 5.3. Thus much use has been made of amphiphiles containing, for example, acryloyl, methacryloyl, vinyl and diacetyl groupS.63 More recently, surfactants containing styrene moieties have also been polymerized.66 In theory, concentrated systems of the polymers produced should result in the formation of liquid crystalline structures. Such materials, however, have not been studied for LC behaviour as their prime interest lies within the field of polymer membranes.

A number of methods have recently been applied to determining the properties of polymer membranes. These include the spreading of monolayers at an air-water interface67 and also the production of vesicles or liposomes.68 The former method allows for the compression of the two-dimensional film between movable barriers. At the same time a film balance registers the change in surface pressure as a function of molecular area, thereby producing a pressure- area isotherm. Various states of molecular orientation can be determined from the shapes of the plot, typified by the diagram in Fig. 5.20. Such states are analogous to those in three-dimensional systems, i.e. solid, liquid, gas.

The polymers are commonly produced by first spreading the film and then inducing the reaction process, rather than by spreading a layer of previously polymerized material. Many studies have been carried out on the variation of reaction kinetics with monolayer state. For example, certain diacetylenic lipids have been shown to undergo polymerization


IT

A -

Fig. 5.20. Surface pressure (n)/area (A) isotherm of a monolayer at the air-water interface: a, gas analogous state; b, liquid analogous state; c, condensed or solid

phase.

only in the tightly packed solid film,69 whereas butadienic or methacryloylic lipids are also polymerizable in the liquid-like phase. 70 The properties of the resultant polymers also vary considerably, depending on their chemical constitution, flexibility and the nature of the polymer linkages. 71, 72 In general, however, the stability of the films to collapse at high surface pressure is significantly enhanced. Monolayer spreading techniques are also useful for studying the effects of varying conditions in the supporting liquid or subphase. Thus the pressure-area isotherms can be studied as a function of temperature, pH and even protein or enzyme concentration.

Liposomes or vesicles are spherical aggregates in which an aqueous interior is separated from the surrounding medium by one or more lipid double layers. 73 They are most commonly produced by the ultrasonication of lipid suspensions, though alternative techniques such as dialysis are also known. In common with monomeric lipid monolayers, such liposomes are usually not stable over long periods of time. However, UV polymerization of a liposomal suspension yields highly stable entities. 74

An interesting application of polymeric vesicles lies in the process of controlled drug release. 63 It has been proposed that such structures could be used to encapsulate materials which would subsequently be released on the enzymatic breakdown of the membrane at the targeted site. The biological impact of such materials is potentially very great and a great deal of research is being conducted in this area. General applications are still some way off, however, owing mainly to the elaborate and critical


in-vivo release mechanism which would have to be worked out. Incorporation of natural lipids into the structure may also be required to enhance flexibility in the membrane or perhaps to act as sites for functional proteins in surface recognition. Further applications have also been proposed. These include uses as anti-tumour agents69 and intermediates in solar energy conversion and artificial photosynthesis. 7 5 Whatever the drawbacks at this stage, there is no doubt that this is an interesting new field of research which will offer increasingly challenging opportunities in the future and a greater collaboration between biochemists and physical chemists.

REFERENCES

1. Morgan, P.W. Macromolecules, 1977, 10, 1381. 2. Kwolek, S.L., Morgan, P.W. Schaefgen, 1.R. & Gulrich L.w. Macromolecules,

1977, 10, 1390. 3. Friberg S.E., Thundathil, R. & Stoffer, 1.0., Science, 1979,205,607. 4. Paleos, CM. Stassinopoulou, CI. & Malliaris, A., J. Phys. Chem., 1983,87,

251-254. 5. Kabanov, VA, Pure Appl. Chem., 1967, 15, 391. 6. Finkelmann, H., Luhmann B. & Rehage, G., J. Colloid & Polymer Sci., 1982,

260,56-65. 7. Blumstein, A. & Hsu, E.C, in Liquid Crystalline Order in Polymers. Academic

Press, New York, 1978, pp. 105. 8. Finkelmann, H., Ringsdorf, H. & Werdorff, 1.H. Makromol. Chem., 1978,179,

293. 9. Luhmann, B., & Finkelmann, H., Colloid & Polymer Sci., 1987,265,506-516.

10. Gallot, B. & Douy, A., Mol. Cryst. Liq. Cryst., 1987, 153, 367-373. 11. Bader, H., Dorn, K., Hupfer, B. & Ringsdorf, H., Adv. Polymer Sci., 1985,64,

1-62. 12. Gray, G.W. & Winsor, P.A. (Eds), Liquid Crystals and Plastic Crystals. Ellis

Horwood, Chichester, 1974. 13. Performance Chemicals, 1986, 1(3), 39. 14. Tanford, C, The Hydrophohic Effect, 2nd edn. Wiley, New York, 1980. 15. Mukerjee, P. & Mysels, K.J., Critical Micelle Concentrations of Aqueous

Surfactant Systems. National Bureau of Standards, U.S. Dept. of Commerce, 1971.

16. Hartley, G.S., Aqueous Solutions of Paraffin Chain Salts. Hermann et Cie, Paris, 1936.

17. Hayter, 1.B. Ber. Bunsenges. Phys. Chem., 1981,85,887. 18. Debye, P. & Anacker, E.W., J. Phys. Colloid Chem., 1951,55,644. 19. Lindman, B. & Wennerstrom, H., Phys. Rep., 1979,52, 1. 20. Mitchell, D.J. & Ninham, B.W., J. Chem. Soc. Faraday Trans. 2, 1981, 77,

601.


21. McMullen, W.E., Ben-Shaul, A., & Ge1bart, W.M., J. Colloid Interface Sci., 1984, 98, 523.

22. Turro, NJ. & Yekta, A., J. Am. Chern. Soc., 1978, 100, 5951. 23. Persson, B.-O., Drakenberg, T. & Lindman, B., J. Phys. Chern., 1979,83,3011. 24. Luzzati, V., in Biological Membranes, ed. D. Chapman. Academic Press,

London, New York, 1971, vol. 1, p. 1. 25. Tiddy, GJ.T., Phys. Rep., 1980, 57, 1. 26. Fontell, K., Mol. Cryst. Liq. Cryst., 1981,63, 59-82. 27. Larsson, K., J. Phys. Chern., 1989,93, 7304-7314. 28. Gruner, S.M., J. Phys. Chern., 1989,93,7562-7570. 29. Lindblom, G. & Rilfors, L., Biochim. Biophys. Acta, 1989,988,221-256. 30. Lawson, K.D. & Flautt, TJ., J. Am. Chern. Soc., 1967,89, 5489. 31. Boden, N., Radley, K. & Holmes, M.C., Mol. Phys., 1981,42,493. 32. Mitchell, DJ., Tiddy, GJ.T., Waring, L., Bostock, T. & McDonald, M.P., J.

Chern. Soc. Faraday Trans. I, 1983,79,975. 33. Fontell, K., in Liquid Crystals and Plastic Crystals, ed. G.W. Gray & P.A.

Winsor. Ellis Horwood, Chichester, 1974, vol. 2, p. 80. 34. Rosevear, F.B., J. Am. Oil Chern. Soc., 1954,31,628. 35. Lawrence, A.S.C., in Liquid Crystals 2, ed. G.H. Brown. Gordon and Breach,

London, 1969, vol. 1, p. 1. 36. Luhmann, B., Finkelmann, H. & Rehage, G., Die Ange. Makromol. Chern.,

1984,123,217-227. 37. Dickstein, w.H. & Lillya, c.P., Polymer Prep. (Am. Chern. Soc., Diu. Polymer

Chern.), 1987,28(1),290. 38. Boden, N., Bushby, RJ. & Hardy, c., J. Phys. Lett., 1985,46,325. 39. Menger, F., Takeshita, M. & Chow, J.F., J. Am. Chern. Soc., 1981, 103,

5938. 40. Keller-Griffith, R., Ringsdorf, H. & Vierengel, A., Colloid & Polymer Sci.,

1986,264,924-935. 41. Larrabee, C.E. & Sprague, E.D., J. Polymer Sci. Polymer Lett., 1979, 17,

749-751. 42. Thundathil, R., Stoffer, J.O. & Friberg, S.E., J. Polymer Sci. Polymer Chern.

Ed., 1980, 18, 2629-2640. 43. Hamid, S.M. & Sherrington, D.C., Polymer, 1987,28, 332-339. 44. Cundy, C.S., Kingston, H.M. & Lappert, M.F., in Advances in Organometallic

Chemistry, ed. F.G.A. Stone & R. West. Academic Press, New York, 1973, vol. 11, p. 300.

45. Finkelmann, H. & Rehage, G., Makromol-Chem. Rapid Commun., 1980, 1, 31-34.

46. Apfel, M.A., Finkelmann, H., Janini, G.M., Laub, RJ. Luhmann, B.H., Price, A., Roberts, W.L., Shaw, TJ. & Smith, C.A., Anal. Chern., 1985,57,651.

47. Gray, G.W., Lacey, D., Nestor, G. & White, M.S., Makromol. Chern. Rapid Commun., 1986, 7, 71.

48. He, x., Lapp, A. & Herz, J., Makromol. Chern., 1988, 189, 1061-1075. 49. Jahns, E. & Finkelmann, H., Colloid & Polymer Sci., 1987,265,304-311. 50. Hall, C. & Tiddy, GJ.T., In SurJactants in Solution, ed. K.L. Mittal. Plenum

Press, New York, 1989, vol. 8, pp. 9-23. 51. Hall, PJ. Williams, G., Tiddy, GJ.T. & Conroy, J.P., to be published.

272 P.J. Hall and G.l.r. Tiddy

52. Finkelmann, H. & Schafbeutle, M.A., Colloid & Polymer Sci., 1986, 264, 786-790.

53. McBain, lW. & Lee, W.W., Oil and Soap, 1943,20, 17-25. 54. Strauss, u.P. & Gershfield, N.L., J. Phys. Chem., 1954, 58, 747. 55. Sprague, E.D., Duecker, D.C & Larrabee, CE., Jr., J. Am. Chem. Soc" 1981,

103, 6797-6800. 56. Hall, PJ., Williams, G., Maud, lM. & Tiddy, GJ.T., to be published. 57. Pietschmann, N., Brezesinski, G., Tschierske, C, Zaschke, H. & Kuschel, F.,

Liq. Cryst., 1989, 5(6), 1697-1709. 58. Gaspard, S., Hochapfel, A. & Viovy, R., Hebd. Seances. Acad. Sci., Ser. C,

1979, 289, 387. 59. Gallot, R & Marchin, R, Liq. Cryst., 1989,5(6),1719-1727. 60. Gallot, R & Marchin, B., Liq. Cryst., 1989,5(6),1729-1735. 61. Kjellander, R., & Florin, E., J. Chem. Soc. Faraday Trans. 1, 1981,77,2053. 62. Conroy, lP., Hall, C, Leng, CA., Rendall, K., Tiddy, GJ.T. & Walsh, l,

Prog. Colloid & Polymer Sci., 82, 1990, 253. 63. Bader, H., Dorn, K., Hupfer, B. & Ringsdorf, H., Adv. Polymer Sci., 1985,

64, 1. 64. Fendler, lH., Science, 1984, 223, 888. 65. Singer, S. l & Nicolson, G.L., Science, 1972,175, 720. 66. Rolandi, R., Paradiso, R., Xu, S.Q., Palmer, C & Fendler, 1.H. J. Am. Chem.

Soc., 1989, 111, 5233-5239. 67. Gaines, G.L., Insoluble Monolayers at Liquid-Gas Interfaces. Interscience,

New York, 1966. 68. Szoka, F. & Papahadjopoulos, D., Annu. Rev. Biophys. Bioeng., 1980, 9, 467. 69. Gros, L., Ringsdorf, H. & Schupp, H., Angew. Chem. Int. Ed. Engl., 1981,20,

305. 70. Hupfer, B., Ringsdorf, H. & Schupp, H., Makromol. Chem., 1981, 182,247. 71. Folda, T., Gros, L. & Ringsdorf, H., Makromol. Chem. Rapid Commun., 1982,

3, 167. 72. Meller, P., Peters, R. & Ringsdorf, H., Colloid & Polymer Sci., 1989, 267,

97-107. 73. Bangham, A.D., Standish, M.M. & Watkins, lC, Mol. Bioi., 1965, 13,238. 74. Regen, S.L., Singh, A., Oehme, G. & Singh, M., J. Am. Chem. Soc., 1982, 104,

791. 75. Fendler, lH., Membrane Mimetic Chemistry. Wiley-Interscience, New York,

1982.

Chapter 6

lyotropic Main Chain liquid Crystal Polymers

M.G. Northolt and D.J. Sikkema Akzo Research Laboratories. P.O. Box 9300. 6800 SB Arnhem. The

Netherlands

6.1 INTRODUCTION AND SYNTHETIC ASPECTS

Lyotropic main chain liquid crystal polymers (LCPs)~forming liquidcrystalline (LC) solutions owing to the conformation of the polymer main chain rather than due to mesogenic character of substituents on the main chain~were first recognized in the case of poly( y-benzyl-L-glutamate).!

Useful reviews of LCPs listing, inter alia, the array of extended chain type (co)polymers that have been found to exhibit lyotropic behavior appeared recently;2.3 however, polyethers are overlooked. Reviews of LC polyethers,4 and of poly(diacetylenes)5 are available. Lyotropic LC polymers are used mainly for fiber spinning. Polymers that have gained technical importance are the para-aromatic polyamides or aramids poly( p-phenylene terephthalamide) (PpPT A), poly( p-benzamide) (PpBA), the heterocyclic rigid rod polymers PBT and PBO (see below), and the aromatic polyamide-hydrazides. The 50/50 (molar ratio) co polyamide from p-phenylenediamine and 4,4' -diaminodiphenyl ether with terephthalic acid is spun into a high-performance yarn from an isotropic solution. We propose to limit discussion to these technically interesting polymers which, naturally, have been studied more thoroughly than the wide variety of (co)polymers that have not (yet) been carried through to (semi)commercial development. Cellulose and its derivatives might be further candidates for high-performance yarn preparation. 6 •7

273

274 M.G. Northolt and D.J. Sikkema

However, there is not yet semicommercial development of cellulose fibers processed via mesomorphic solutions. The statistical physics of lyotropic LCPs has been dealt with in Ref. 8, but a modified Maier-Saupe theory developed by Picken will be discussed in this chapter.

PpBA

PpPTA

PpBAT

6.1.1 Aromatic Polyamides Recent in-depth reviews dealing with PpT A and PpBA and a review covering a wide range of aromatic (co)polyamides are available.9 ,lo

Synthesis of PpBA requires the synthesis of the hydrochloride salt of 4-aminobenzoyl chloride; this is achieved by careful treatment of Cl-CO-C6H4 - N=S=O with HCI,u The sulfinylaminobenzoyl chloride is prepared by treating 4-aminobenzoic acid with excess thionyl chloride. Polymerization is achieved by simply dissolving the monomer hydrochloride salt in an amide solvent. The resulting molecular weights remained relatively modest. Alternative methods of preparation did not lead to improvements. 9 Much more work has been carried out with respect to the preparation of PpPT A. The monomers l,4-phenylene diamine and terephthaloyl chloride are commercially available. The commercially successful polymerization employs these monomers in a complex solvent consisting of N-methylpyrrolidone and calcium chloride. The discovery that more CaCl2 should be used than the 6 wt% that can dissolve in NMP at room temperature (the polyamide formed dissolves

Lyotropic Main Chain Liquid Crystal Polymers 275

further CaClz; the temperature of the mixture rises upon the start of the exothermic polymerization) was crucial in arriving at the optimum process. 1Z The historical development (an earlier process for PpPT A used a mixture of NMP and the carcinogenic hexamethyl phosphoric triamide) has been described in detail. 9

When aiming at a final polymer concentration of 10 wt%, and 11 wt% of CaCh relative to NMP, the CaClz is first slurried with dry NMP; the p-phenylenediamine is dissolved in this slurry by heating to about 80°C-or by adding molten PPO without further external heating; this mixture is cooled to O°c. Then the terephthaloyl chloride is added quickly (either as flakes or as a melt) to the PPO-NMP-CaClz slurry, which is stirred at a high rate. Immediately the viscosity of the reaction mixture rises to a very high level. Stirrring/kneading is continued forcefully; soon the stirrer torque decreases as the mixture breaks up into crumbs. Kneading is continued for another 10 min before the product is dumped into excess water to wash out NMP, HCl and salt. It may be desirable (for corrosion reasons) to add Ca(OH)z to the water. The molecular weight of the PpPT A can be controlled by deviating from the exact stoichiometry of PPO vs TOC, or by allowing the presence of a small amount of water. In the exactly similar preparation of PpBAT (from 4,4'-diaminobenzanilide and TOC) the presence of traces of water has a much smaller influence on the final degree of polymerization. The polymerization kinetics were investigated in model systems 13 and have been re-investigated. 14 This very fast process has-in the absence of limitations from the rapidly rising viscosity-a rate constant of about 1000 litre mol- 1 s -1, an activation enthalpy LlH# of about 30 kJ mol- 1

and an activation entropy LlS# of about - 80 J mol- 1 K -I. Alternative processes for PpPT A have received wide attention but

usually fail to lead to polymer of sufficiently high molecular weight. A few reports on PpPT A with an acceptable degree of polymerization, prepared via unconventional routes exist: the Yamazaki phosphite-pyridine condensation agent was coaxed to produce [1J]=4'5dlg- 115 and even 6·2dlg- I . 16 ,17 A vapor-phase reaction of PPO and TOC gave inherent viscosities as high as 5·3 dl g - 1.18 N -silylated PPO was polymerized with TOC to high-d.p. PpTA. 19 Finally, PpPTA almost of sufficient molecular weight was reported by reacting PPO'2HCl with terephthalic acid in NMP-CaCIz with SOClz.zO

Aromatic copolyamides, including the 50/50 (molar ratio) copolymer of 3,4'-diaminodiphenyl ether with terephthaloyl chloride, are synthesized in a similar manner using amide solvent-salt combinations-in practice


NMP-CaCI2 . This 50/50 copolymer is in fact not processed via lyotropic solution, but rather directly as the (relatively dilute) polymerization reaction mixture; the impressive yarn properties are realized only after hot drawing. Concentrated LC solutions in sulfuric acid are inaccessible owing to degradation of the polymer in that medium.

Aromatic hydrazides and amide-hydrazides can similarly be prepared by condensation of diacid chlorides and hydrazine or preformed aromatic hydrazides in amide solvents, usually with salt added. Most attention in this field has been centered on the adduct from 4-aminobenzohydrazide and terephthaloyl chloride; this product was not processed via a liquidcrystalline solution, although lyotropic behavior under special circumstances was reported: processing was carried out using the amide-salt solution resulting from the polymerization; LC behavior was observed in sulfuric acid solution.2.3

6.1.2 Rigid Rod Heterocyclic (Ladder) Polymers Among the wide variety of ladder or ladder-like polymers described in the literature, which are aimed at enhancing the thermal stability rather than at superior fiber-mechanical properties,21,22,37 the benzazole polymers of the so-called 'PBZ' family have attained a semicommercial status as high-performance fibers processed via a liquid crystalline solution. A series of papers describing various aspects of these polymers appeared some time ago,23-30 and an excellent and comprehensive review appeared recently.31

The benzoxazole and the benzothiazole polymers which received most attention are the cis and trans isomers, respectively. PBT, which has been the object of most of the work, is prepared by first synthesizing 2,5-dimercapto-l,4-phenylenediamine: PPD is converted into the bisthiourea with ammonium thiocyanate. This product is cyclized with bromine and hydrolyzed to the oxidatively exceedingly fragile


monomer potassium salt, which is isolated under reducing conditions with HC1.24

S II

H2NCHN~ • i NH,SCN i~~

NHCNH2

I glacial

8r2 acetic i aCid ,

II S

The bis-HCI salt is easily oxidized also, but can be handled. The tetrafunctional monomer for PBO is prepared from resorcinol diacetate by nitration and hydrogenation. Care must be taken in view of explosive side products. A detailed description appeared recently.32 A probably more economical synthesis of 4,6-diaminoresorcinol has become available by nitration of 1,2,3-trichlorobenzene followed by alkaline hydrolysis and hydrogenation.33 The status of the work aimed at a commercial process for cis-PBO was reviewed recently.34 Diaminohydroquinone, the prec~rsor for trans-PBO, can be prepared from chloranil by ammonolysis followed by hydrogenation: 35

CI CI

CI'©rCI ClkCI o HN03 • JQl NaOH.

~N N02

Polymerization is brought about by dissolving the tetrafunctional monomer in polyphosphoric acid (PPA), by slowly heating to about

278 M.G. Northolt and D.l. Sikkema

lOODC and evacuating in order to eliminate the hydrochloric acid. This must be done in PPA of limited strength to avoid problems with foaming due to too high a viscosity of the PPA. After the HCl has been removed, micronized terephthalic acid is added and sufficient P 20 5 to bind the water of polycondensation and ring closure, resulting in a solution containing between 82 and 84 wt% of P20 5 in the PPA. The mixture is stirred and slowly heated to 180"C over several hours. The rate of polymerization appears to increase after the mixture becomes liquidcrystalline. 36 The resulting polymer solution can be spun.24.31

An alternative route to polybenzazoles was devised, by polymerizing dinitro difunctional nucleophilic monomers with terephthaloyl chloride, followed by reduction of the nitro groups and cyclization. 37 This approach has the advantage of oxidative stability of the dinitro monomers; complete conversion in reduction and cyclization are unlikely to be easily achieved.

Finally, we mention the interesting exploratory work on the structurally modified PBZ polymers which has been carried out by the U.S. Air Force. 38 40

6.2 ORDER IN LYOTROPIC POLYMER SOLUTIONS

The development of liquid-crystalline phases and the rheological properties of lyotropic polymers are discussed in so far as they are related to the formation of fibers and films.

The mechanical properties of fibers and films are governed by the orientation distribution of the polymer chains (see Section 6.4.1). Hence the discussion is focused on the development of the orientation of the chains in the polymer solution and on the formation of orientation during spinning of fibers. Good surveys of current theories on the formation of liquid-crystalline phases have been given by Odijk 8 and Ciferri and Marsano. 41 Later in this chapter we shall present a modified Maier-Saupe theory recently developed by Picken of our laboratory.42 45

In order to form liquid crystals, the polymer chain should have sufficient intrinsic rigidity. Stiff-chain polymers are defined as polymers with a Mark-Houwink exponent ([1]] = K'M") greater than lover some range in molecular weight. Because rigid polymer chains are almost insoluble, strong solvents are needed to dissolve these polymers. The rheological behavior of the lyotropic solutions of the technologically important aromatic polyamides has been studied extensively. Recent


work by Picken et al. deals in particular with the structure and properties of a nematic solution under shear and elongational f1oW.ll.42-54

Poly(l,4-benzamide), abbreviated here as PpBA, was the first nonpeptide, synthetic condensation polymer reported to form liquidcrystalline phases. I 1.46.49,55-57 Of technological importance is poly(l,4-phenylene terephthalamide), abbreviated here as PpPT A. In both polymers the non-flexibility or semi-rigid chain structure is caused by the para-linked benzene rings and the partial double bond character of the carbon-nitrogen bond in the predominantly trans amide linkages.

PpBA and PpPT A can form liquid-crystalline solutions in selected N,N-dialkylamides with or without added lithium or calcium chloride, or in strong acids, such as sulfuric acid and chlorosulfonic acid.!! The nematic phase in a quiescent solution of these polymers starts to develop at a critical concentration, c*. This phase consists of domains with more or less parallel-oriented chains, and the amount extends rapidly with increasing concentration. In general, the critical concentration is a function of the temperature, molecular weight and molecular weight distribution, chain rigidity and the solvent applied. The liquid-crystalline solutions of aromatic polyamides in the quiescent state are turbid in appearance and optically birefringent, but become opalescent owing to shearing and stirring.

In the range below the critical concentration the viscosity of the solution increases with increasing concentration, but at the critical concentration this viscosity reaches a maximum and with the onset of mesophase formation it drops rapidly to a minimum. As the concentration is further increased, the viscosity rises again until the point of solidification. From experiments on a series of PpBA solutions of different molecular weight, Papkov derived general concentration dependences of the viscosity by reducing the concentration to c* and the viscosity to the maximum viscosity at the phase transition.46 A theoretical treatment of the behavior of the viscosity near the critical concentration has been given by Kawai 58 and Doi. 59 The anisotropic phase has a higher polymer concentration than the coexisting isotropic phase and is composed of the higher-molecular-weight chains. I I Furthermore, the critical concentration increases as the molecular weight of the polymer decreases. A hypothesis on the solution structure of the aromatic polyamide chain, in particular the specificity for fully substituted amides as effective solvents and the function of salts in mixed solvents of this type, has been made by Panar and Beste.49


Characterization of dilute solutions of PpPT A and PpBA by determination of the persistence length, Lp , and the molecular weight distribution has been carried out by Arpin and Strazielle,48 Ying and Chu,60 and Ying et al. 6 1.62 The results depend on the experimental method, such as the determination of the radius of gyration, the rotary diffusion constant from flow birefringence or the optical anisotropy from depolarized light scattering. Furthermore, the solvent affects the rigidity of the chain, and thus Lp. For PpBA the reported Lp values are higher than for PpPTA, viz. about 50 nm against 30 nm. A satisfactory explanation for this difference has not yet been given. The relation between viscosity and molecular weight of PpBA and PpPTA in solutions of N,Ndimethylacetamide 4 wt% LiCI and chlorosulfonic acid has been studied by Schaefgen et al. using light scattering.63 For both polymers they found that the exponent a in the Mark-Houwink relation is 1·1 for M w > 12000 and 1·7 for M w < 12000. The latter value for the exponent was also found by Tsvetkov et al. for PpBA dissolved in N,N-dimethylacetamide, but in the range M w < 17000.64 Other studies using gel permeation chromatography68 and laser light scattering61 yielded a = 1·2 for M w < 25000, while for M w > 25 000 the exponent became smaller. The decrease of the Mark-Houwink exponent for ranges with larger molecular weights is typical of worm-like chains.

Viscosity measurements of solutions of low-molecular-weight PBO in chlorosulfonic acid and other solvents yielded a Mark-Houwink exponent of 1·85, which is very close to the value expected for a rod-like polymer. 52 Depolarized dynamic light scattering of PBT solutions in chlorosulfonic acid resulted in a persistence length of 64 nm, which is one of the highest persistence lengths known when compared with those of other rigid-rod polymers.65

The role of electrostatic interactions and molecular association on the rheological properties of PpPT A and PBO solutions has been studied extensively by Berry and co-workers. 52-54 For dilute solutions in a low-ionic-strength acid, an extremely small translational diffusion coefficient and a high viscosity have been observed, which they attributed to a pseudo-ordering of the polymer solvent system caused by electrostatic repulsions between the protonated polymer chains. Extrapolation of the data at infinite dilution gave an Lp of 45 nm for PpPT A, which is larger than previous results. Failure to recognize the effect of electrostatic repulsion can lead to erroneously low estimates of the persistence length.

A measure of the average orientation in a domain of a nematic liquid is the order parameter (P2)' This is the average of the second-order


Legendre polynomial of (cos q», where q> is the disorientation angle between the chain axis and the director. It can range from 0 for random orientation to 1 for perfect parallel orientation. Order parameters for nematic solutions of PpBA in N,N-dimethylacetamide+3% (w/w) LiCI were determined in the composition range extending from the lower limit for stability of the pure meso phase, which is just below the critical concentration, to the solubility limit of the polymer.67 The experimental values of the order parameter, ranging from 0·76 to 0'83, were found to lie between the theoretical predictions of Doi 59 and those of Flory and Ronca. 68

The structure of the solid state of nematic aramid solutions will be discussed in Section 6.3.4.

6.2.1 The Modified Maier-Saupe Mean Field Theory Two mechanisms can lead to the formation of a nematic phase in a liquid solution. The first is based on the traditional Onsager or Flory type of approach. 69. 70 This treatment predicts that at a certain concentration the molecular asymmetry alone is sufficient to create an ordered phase without any attractive interactions. The second approach shows that a stable meso phase is formed by an anisotropic potential.

The excluded-volume theories of Onsager and Flory show that, depending on the axial ratio of the rod-like particles, there is a critical concentration above which a nematic phase is formed. This concentration does not depend on the temperature of the system as these theories are essentially athermal, i.e., the part of the free energy leading to the anisotropy is an entropy term.

A later version of the Flory theory describes a system of semi flexible particles. The persistence length is used to determine an effective axial ratio for the particles. 71 A temperature-dependent persistence length thus leads to a 'thermotropic' type of behavior where the phase transition is governed by temperature and concentration.

The second or alternative mechanism is used by the Maier~Saupe mean field theory in which the stability of the nematic phase is derived from an anisotropic potential. Picken has developed a theory for the nematic phase formation of Iiquid~crystalline polymers which is based on the Maier~Saupe mean field theory.42.43 A molecule in a nematic domain, with its axis at an angle q> with respect to the director or average orientation axis of the domain, is assumed to feel the influence of the surrounding medium only in terms of an anisotropic potential

(6.1)


where q is a constant describing the strength of the orienting potential and <P2) is the value of the second-order Legendre polynomial, P2(cos cp)=t (3 cos2 cp-1), averaged over the orientation distribution function f(cp) of the molecular axes. Using this potential we obtain for the orientation distribution

where k is the Boltzmann constant, Z is the partition function

Z = Il exp [k~ <P2) P2(cos CP)] d(cos cp)

and <P2 ) is found from the equation

(6.2)

(6.3)

<P2)=Z-1 Il exp[k~<P2)P2(Coscp)]p2(COSCP)d(COSCP) (6.4)

On equating the free energy in the isotropic and anisotropic phase this equation leads to a first-order transition at kT/q = 0·22. The order parameter <P2) at this transition is 0-43. Given kT/q, <P2 ) can be derived from eqn (6.4) and f(cp) is found from eqns (6.2) and (6.3).

Picken has shown that the model of a worm-like chain seems to be the best to describe the properties of aromatic poly ami des dissolved in a good solvent such as sulfuric acid, in particular the temperature dependence of the persistence length. Hence, for an adaption of the Maier-Saupe theory to lyotropic polymer solutions the influence of the concentration and of the flexibility, i.e. the deviation from rigidity, has to be taken into account. The latter is demonstrated by the fact that the persistence length of PpPT A in a sulfuric acid solution is 29 ± 5 nm,60 whereas the contour length for a characteristic molecular weight, M w, of 30000 is 170 nm, which implies that the molecules are far from being completely rigid.

The Maier-Saupe theory is now adapted to take concentration and flexibility into account via the strength of the anisotropic potential. This is done by power-law relations. First the 'contour projection length' L(T) is introduced. This is defined as the projection of a polymer chain along the direction of the first segment. Next it is assumed that the strength of the orienting potential is given by

(6.5)


where q* is a scaling constant and c is the concentration. The dependence on the factor c2 is based on the fact that the attractive part of the Lennard-lones potential is proportional to r- 6, where r is the distance between the molecules. This dependence is proportional to V - 2 or ('2,

where V is the volume. The Maier-Saupe potential is a generalized two-particle interaction for which an L 2 dependence may be expected. Using a worm-like chain the following form for the contour projection length can be derived: .

[ LeT] l-exp -L T L(T)=Lp T p p (6.6)

Tp

where Lp is the persistence length determined at an arbitrary temperature Tp and Le is the contour length. The limiting value for an infinite chain is Lp, whereas the value for a very short molecule is the end-to-end distance. Thus L(T) describes the temperature dependence of the rigidity of a semiftexible polymer for all molecular weights.

At the nematic-isotropic transition temperature Tn;, also called the clearing point, L(Tn;) is determined by eqn (6.6). Thus, by measuring Tn; as a function of the concentration c, the scaling constant q* can be determined from the relation

(6.7)

which can also be written as

(6.8)

It can be shown that for low-molecular-weight polymers IX is larger than for high-molecular-weight polymers.

The anisotropic potential, U, is now known for other concentrations, temperatures and molecular weights, so eqn (6.4) can be used to calculate the order parameter (P2> as a function of the temperature.

Figure 6.1 shows the clearing temperatures as a function of the concentration observed for PpPT A solutions in sulfuric acid, while Table 6.1 summarizes the results obtained for PpBA, PpPTA and PpBAT (poly(4,4'-benzanilidylene terephthalamide)). Good agreement with eqn (6.8) is obtained. The slope of the concentration--clearing temperature curves depends on the molecular weight, which indicates coupling of the


180.,------------------,

140

-100 ~

~ 60

20

*

• * •

• Mw=12000 • * Mw=31 000

-20+-~~-~~-~-~~~~ 7 8 9 10 11 12 13 14

Concentration (%w/w)

Fig. 6.1 Clearing temperature Tni as a function of polymer concentration (w/w) of PpPTA in 99·8% sulfuric acid. The average molecular weights Mw are indicated by different symbols. The drawn curves are from eqn. (6.7) and q* is

determined at the filled circle.42

Table 6.1 Results of a least-squares fit of the clearing temperature (Tni) versus concentration (e) measurements, to relation (6.8) for various aromatic aramid solutions. The standard deviations of A and (1 are shown in

parentheses

Sample Mw A C(

PpBAT 8000 37(2) 0·92(3) PpBAT 42000 74(5) 0·66(3) PpBAT 70000 115(10) 0·49(4) PpPTA 12000 46(4) 0·85(4) PpPTA 31000 99(2) 0·56(1) PpBA 10000 41(6) 0·92(6)

(temperature-dependent) molecular flexibility to the phase transition as expressed in eqn (6.6).

The anisotropy of the polarizability, de, is a good estimate of the order parameter in a nematic solution:

(6.9)

Details of the experimental method are given in Ref. 72. Figure 6.2 shows the anisotropy of the polarizability (at optical frequencies) as a function of the temperature measured for an anisotropic PpBA T solution. The solid curve results from Picken's theory, whereas the dashed curve is derived from the standard Maier-Saupe model.

Lyotropic Main Chain Liquid Crystal Polymers

0.2

0.1

r---------------.to

._.-.-.-._.-.-._.-., ..... ...... " 05 ;p

V

OO+----r----,~---,---+_-~-+O.O -60 -40 -20 0 20

T-Tni (K)

285

Fig. 6.2. Anisotropy of the dielectric constant and <P2), as a function of the relative temperature (T - TnJ for M w = 8000. The drawn curve is from Picken's theory. The dashed curve is from the traditional Maier-Saupe

model.

Although the influence of the excluded volume is not taken into account, Picken's theory provides a satisfactory interpretation of the experimental results. The combination of a worm-like chain model and the Maier-Saupe approach has also been suggested by Jiihnig,73 and is developed further by ten Bosch, Maissa and Sixou,74 as well as by Warner et ai. 7 5

These theoretical considerations have led to the following view of a nematic solution, in particular of a solution of a para-aromatic polyamide in sulfuric acid.43 In a quiescent solution of a lyotropic polymer the chains are aligned more or less parallel inside domains of microscopic size (see Fig. 6.3). The degree of orientation inside the domain, as represented by the order parameter (P2), is determined by the concentration and temperature. The excluded-volume entropy term leads to the formation of 'oriented blobs' with a size of the order of Lp, the persistence length. These blobs line up owing to their anisotropic polarizability, which implies that the formation of the anisotropic phase is governed by a dipole-dipole type of interaction, immediately leading to the Maier-Saupe mean field potential. The entropy or excludedvolume interaction merely tells us something about the local molecular structure and not about the long-range orientational order. On a macroscopic scale a random director field is found for anisotropic samples in an equilibrium situation, i.e. if the samples have relaxed for a sufficiently long time after the last deformation and there are no external orienting effects.


ii

Lp

Fig. 6.3. Schematic structure of a single domain in an aramid solution. Each domain contains several 'blobs' of dimension Lp. The director n is the average

orientation of the molecules inside the domain.

6.2.2 Flow Behavior of Lyotropic Solutions The structural transformations occurring near the critical concentration cause changes in the flow properties of the solutions. In dilute solutions (c < c*) one observes a region of Newtonian flow at low shear rates and a region of shear thinning at higher shear rates. For c ~ c* the rheological behavior is very different. There is also a region of quasi-Newtonian flow, but as the shear rate decreases the apparent viscosity starts to increase rapidly, indicating plastic behavior. At high shear rates the maximum in the viscosity versus concentration curve disappears as is shown in Fig. 6.4. This phenomenon may be explained assuming that at high shear rates very little difference exists between the anisotropic and isotropic phases,46,54, 76, 77 as illustrated by birefringence measurements by Picken84 presented in Fig. 6.5.

Rheological properties under steady state and oscillatory shear flow of isotropic and nematic solutions of PpPT A, PBT and PBO were studied by Baird 77 and Berry et al. 5 3 Baird observed shear thinning for a series of PpPTA solutions in sulfuric acid (4-15%). These results also suggest that at higher shear rates very little difference exists between the anisotropic and isotropic phases. Steady-state viscosities as a function of the temperature observed for solutions of PBO in methanesulfonic acid


~1000 Q)

'" o Cl

.?: iii o ~ 500 :;

Shear rate

-100 sec-1

*300sec-1

.. 500sec-1

O+-~~-r-r~~-r-.~~-.~

o 5 10 15 Concentration (wt%)

287

Fig. 6.4. Viscosity of PpPT A solutions with l7inh = 4·2 in concentrated sulfuric acid at different shear rates at 40°C.

008

0.06

c

<I 004

0.02

55 65 T(OC)

75 85

Fig. 6.5. Shear-flow-induced birefringence as a function of temperature and shear rate for an isotropic 9% (wjw) PpPTA solution in sulfuric acid.

showed a sharp increase near Tnj , a behavior which has also been reported for PpPT A and PpBA solutions. 53

When the viscosity is plotted against the shear rate on double-logarithmic axes, three regions have generally been observed: at very low shear rates a shear thinning region; next a Newtonian plateau; and at high shear rates a shear thinning region again. An interpretation of the rheology of liquid-crystalline polymers based on this three-region curve has been given by Onogi and Asada 78 and Wissbrun. 79 Detailed rheooptical studies on PpPT A solutions in sulfuric acid were carried out by Onogi, White and Fellers.80•8 ! In the biphasic region (9-9'5 wt%) the PpPT A solutions at equilibrium display small globular structures with


sizes of at least 5 .urn. They have a negative birefringence, indicating a tangential orientation of the chains in these spherulite-like structures. At slightly higher concentrations aggregates of the globules of about 50.um form, which are separated by disclination lines. Under shear flow the isotropic solutions show a steady-state birefringence that increases linearly with the shear rate, and near the isotropic-to-nematic state transition there is a large increase in birefringence.82

Panar et al. describe the effects of a strong magnetic field on a nematic solution of low-molecular-weight PpBA. A transparent state is created in which the domains have grown to a size that is too large for light scattering. Slight stirring will cause the solution to break up into domains in the micrometer range. Shear of a thin layer of a nematic solution of PpBA in an amide solvent produced a pattern of parallel bands of approximately l.um width, in which the molecules are arranged alternately in two directions at 45° to the shear direction.49 This banded texture appears to be a characteristic phenomenon of lyotropic and thermotropic polymers. Marrucci83 has attempted to explain this texture and proposed that it is a manifestation of ordered tumbling of the domains. See Section 6.3.4 for a discussion of the banded texture.

Shear flow-induced birefringence measurements of an isotropic solution of PpPT A in concentrated sulfuric acid with a clearing point of 45°C were also performed by Picken.84 As shown in Fig. 6.5, the flow-induced birefringence increases strongly when the isotropic-nematic transition is approached. The results demonstrate that the application of a relatively small shear rate already leads to a degree of orientational order in the initially isotropic solution that is comparable with the order in the nematic phase. This points to a strong coupling between the orientation and the external flow field, and to the occurrence of a shear-induced phase transition.

The texture during flow shows a rather grainy structure giving the impression of small domains tumbling over each other. The apparent size of the domains decreases with increasing shear rate. After cessation of flow, a banded texture is formed with the bands oriented normal to the shear direction.

In a molecular theory describing the dynamics of rod-like polymers in concentrated solutions, Doi predicted an increase of the order parameter with a shear flow gradient. 59 There have been several theoretical studies which predict that an elongational flow gradient also increases the order parameter substantially. Marucci and Ciferri found, however, that the largest contribution to the order parameter should arise from an increase


of the polymer concentration, but that the effect of an elongational flow field should be more pronounced for less-concentrated solutions and for less-rigid polymers.8s

Results of transient shear-stress measurements on a 20% (w jw) nematic PpPT A solution in sulfuric acid are shown in Fig. 6.6, where the transient shear stress normalized to the steady-state value is given versus the total amount of applied shear.86 A power-law relation was observed between the steady-state shear stress and the shear rate. For the range of the applied shear rate the transient behavior, i.e. the positions and heights of the minima and maxima, is only related to the total shear strain that is applied. An explanation is that in the beginning of the shear flow there are relatively few disc1inations per unit volume, leading to relatively large domains in the nematic solution. At the same time a high degree of orientation is obtained, leading to a low viscosity in the first minimum. Applying further shear leads to the grinding of large domains into smaller ones, implying an increase of the disc1ination density. Finally, the orientation induced by the shear flow and the disc1ination density reach an equilibrium corresponding to a steady-state value of the shear stress. This qualitative model also explains the observation that the relaxation time required to restore the initial structure depends on the total shear strain that is applied. 86

In two supplementary studies applying X-ray and light-scattering techniques it was found that the transient behavior is governed by the

1.5,----------------_-, Temp80'C

---- l' -15.1

-- l' -25.1

- 1'-45.'

10

0.5

O+---~r_--~----~----~--_,,_--~ o 100 200 300

pI 400 500 600

Fig. 6.6 Transient shear-stress measurements presented in dimensionless variables for a 20% (w /w) PpPT A solution in sulfuric acid.


defect density or disclination density.44.45 However, the disclinations hardly affect the average degree of director orientation of the solution, implying that the volume fraction of the disclinations is small. Furthermore, the course of the average director orientation observed during the start-up of a shear flow is explained using an affine deformation model. As will be shown in the next section, this model can also be applied to the elongational flow field which is present during spinning of nematic solutions.

The effect of an elongational flow field on a nematic solution is, according to Panar, rather similar to the observed effects of a magnetic field. 49 Calculations by Khokhlov and Semenov, and Maissa et al., show that an external field does shift the isotropic-nematic transition to lower polymer concentrations and that the biphasic gap disappears at sufficiently high strain rates. 87.88 According to calculations by Bahar and Erman, the effect of an elongational flow on a system of rod-like particles is pronounced when the system is close to or in the biphasic state, but the effect is of minor importance when the quiescent solution is wholly nematic.89 However, no experimental proof of these predictions is yet available.

6.2.3 Spinning of Lyotropic Solutions In this section the development of orientation in the spinning process is discussed, mainly based on the results of experiments performed in our laboratory (see Figs 6.7-6.10). The actual fiber formation by coagulation will be given attention in Section 6.3.4. A useful survey of the various factors involved in the solution spinning of rigid and semi-rigid polymers has been given by Ciferri and Valenti.90

During spinning an external force is applied, which in the spinneret has the effect of an elongational flow in conjunction with a shear flow, but causes an elongational flow in the spin-stretch phase of the spinning process. This elongational flow field will orient the directors of the domains in the nematic solution along the direction of the spinline, whereby increasing draw ratios yield a higher degree of director orientation. Hence the draw ratio, in conjunction with the polymer concentration of the solution, determines the final orientation distribution in the spinning dope before removal of the solvent, as is clearly demonstrated by experiments of Weyland. 76 If the relaxation time of the oriented spinning solution is long enough, the order parameter at the end of the spin-stretch stage can be preserved by solidification and/or coagulation in the spinning bath.


After spinning, the removal of the solvent in solutions of lyotropic polymers is carried out either by evaporation of the solvent (dry spinning) or by coagulation in a non-solvent (wet spinning). In the case of spinning of PpPT A solutions it was found that when the spinneret was placed in a coagulation bath of water, draw ratios (A) not larger than 2 could be obtained for the whole concentration range, as shown in Fig. 6.7. This is

10

8 r"'-'" 6 \ '" x

j . E

~ 4

'" "-", '" 2 -.-._.::=i=::::: ... _ ..... _._

0 0 5 10 15

% cone.( w/w) PpPT A

(a)

70

60

50 E Z 9 40

<IJ ~ ::J

-0 30 0 ~

20

. . 10

0 0 5 10 15

% conc.(w/w) PpPTA

(b)

Fig. 6.7. (a) Maximum draw ratio for spinning as a function of the PpPT A concentration in sulfuric acid. The critical concentration is near 8%. Temperature of the spinning solution and the coagulation bath are both 30De. (D.) Air-gap spinning; (.6) spinneret placed in coagulation bath. (b) Modulus of the fibers spun

according to the conditions in (a).


caused by rapid coagulation of the solution. 76 Higher draw ratios are usually obtained either through a non-coagulating first spinning bath or by placing the spinneret just above the bath, which is then called 'air-gap' spinning.91 The air gap has previously been applied in the spinning of solutions of cuproammonium cellulose.92

Figure 6.7 also shows the large effect of the nematic phase formation on the 'spinnability' of the polymer solution. Below the critical concentration, c*, the maximum draw ratio, A rna'" increases rapidly with decreasing concentration of the solution. Above c* a rapid increase of Ie max is observed for increasing concentration. For a spinneret of 60,um diameter, a draw ratio of 4 and a spinning speed of 100 m min -1, the elongation rate is of the order of 100 s - 1, which is much lower than the shear rates in the spinneret. Therefore, the decrease of the maximum in the shear viscosity at the critical concentration observed for increasing shear rates, as shown in Fig. 6.4, is presumably of no significance to the air-gap spinning process for lyotropic polymers. Results of dynamic shear measurements at low frequencies, however, provide an explanation of the spinning results with solutions having a concentration near c*. Figure 6.8 shows data obtained at 0·025 Hz and 35°C for a series of PpPTA solutions in sulfuric acid. At the critical concentration the behavior of the solution is highly elastic (In(tan b) is small) and the shear viscosity reaches

G"

G'

tanS

Concentration 1% (w/w)

Fig. 6.8. Shear moduli, G' and G", and tan (j obtained from dynamic shear measurements on a series of PpPT A solutions in sulfuric acid at 0·025 Hz and

35°C.


a maximum. Apparently it is this elastic behavior of the solution that limits the draw ratio in air-gap spinning at concentrations near this maximum.

Conio et al. studied the fiber formation of PpBA from an organic solvent, N,N'-dimethylacetamide containing 3% LiCI, but in these experiments an air gap was hardly used.93 This may be the reason why they arrived at somewhat different conclusions concerning the spinning of para-aromatic polyamide solutions.

The effect of the draw ratio on the orientation development during spinning has not yet been extensively reported. Using the simplifying assumption that the alignment of the domains in the solution from elongational flow is similar to the case of dilute suspensions of rigid-rod particles, Kenig arrived at the equation

(6.10)

where CPo is the angle of the rod with respect to the direction of the spinline immediately after the spinneret, cp is the angle at the end of the air gap, ), is the draw ratio and p is the so-called the 'orientability' parameter. Four thermotropic systems and one lyotropic system (PpPTA in concentrated sulfuric acid) were studied. The values found for p ranged from 0·2 to 0'6, the latter being observed for the PpPTA solution.94

Figure 6.9 presents the modulus of PpPT A fibers as a function of the applied draw ratio for air-gap spinning at a very low winding tension.

m,------------------------------------,

60

~* ;::-E fl* * * z

S2 * /.** "' 50 .2 !f ::J "0 0 :2

40

30 1 3 5 7 9 11 13 15

Draw ratio A,

Fig. 6.9. Modulus of PpPTA fibers versus draw ratio in the air gap for a polymer viscosity of '1inh =4·5 and a concentration of the sulfuric acid solution of

19·3%. Estimated upper limit of the winding tension is 7,2 MN m - 2,95


This figure demonstrates the importance of the spin-stretch stage in the spinning process. It will be shown here that these results are in agreement with the affine deformation mode1.95

For the affine deformation, which describes the orientation of a rod-like particle embedded in an elastic matrix, the parameter p in eqn (6.10) equals 3/2. Kuhn and Griin,96 as well as Ward,97 derived a relation between the birefringence and the draw ratio which, in fact, is the relation between the order parameter <P2) and the draw ratio:

1 [3 3K cos -1 K ] <P2)="2 I-K2- (l_K 2)3/2- 1 (6.11 )

where K=)_-3/2 and <P2 )=1-1 <sin 2 cp). It is easily shown that, at least for }_ > 5, eqn (6.11) yields

(6.12)

The overall order parameter, <P2 ), for an oriented nematic solution may be expressed in terms of the order parameter of the director field of the domains, P 2, and of the molecular order parameter, <P2 ), characterizing the degree of order inside the domains, according to

(6.13)

It is now postulated that the affine deformation acts only upon the director field of the solution, according to eqn (6.11) in which <P 2) is thus replaced by P 2' On the assumption that coagulation does not effect the orientation distribution of the chain axes, the overall order parameter <P2 ) can be obtained from the modulus of the fiber (see eqn (6.15) in Section 6.4.1). Knowing the polymer concentration of the spinning solution, the molecular weight of the polymer, the persistence length and the coagulation temperature, one can calculate the molecular order parameter <P2 ) using Picken's modified Maier-Saupe theory. The order parameter of the directors of the domains created by the spinning process is now derived from eqn. (6.13).

Before entering the air gap the spinning solution is subjected to a so-called pre-draw, Ao, due to the spinneret. Hence, the effective draw ratio pertaining to the affine model, A, is given by

A=AOA.

where }_. is the draw ratio in the air gap.

(6.14)


Figure 6.10 presents the curve for the affine deformation model calculated with eqn (6.11) in terms of log (sin2 cp) versus 10gA for the range 8 < A < 80. The experimental values for the overall order parameter

(P2 ) were derived from the modulus values presented in Fig. 6.9. Division of these values by the molecular order parameter (P2 ) =0'925, corresponding to E = 60 GN m - 2, yielded values for the order parameter of the director field P 2, and so the corresponding experimental values of (sin2 cp) shown in Fig. 6.10. By assuming a pre-draw ratio Ao = 8, the observed data plotted against the effective draw ratio were found to coincide with the theoretical curve of the affine deformation. This pre-draw value is in agreement with the one derived from the dimensions of the entrance cone of the spinneret used in these experiments. Note that an incorrect value for (P 2) has a considerable effect on the experimental (sin2 cp) values derived for large draw ratios, thereby causing a change of slope of the curve determined by the experimental data in Fig. 6.10. The curve for the affine deformation has also been drawn in Fig. 6.9.

0

-1

II "s-c: -2 '00 V Ol .2

* -3

-4 0.9 1.1 1.3 1.5

log A 1.7 1.9

Fig. 6.10. Degree of orientation given as (sin2 qJ> versus the effective draw ratio for the range 8 < A < 80. The drawn curve presents the affine deformation model (see text). The experimental data have been derived from the modulus values

given in Fig. 6.9.95

The final degree of orientation attained during the elongational flow stage in the spinning process of a nematic solution cannot be more higher than the molecular alignment inside the individual domains, which is given by (P2> and determined by the thermal fluctuation and the


concentration of the solution. This explains the leveling of the modulus in Fig. 6.9 at high draw ratios.

In conclusion, the work of Picken et al. has provided a theoretical framework supported by many experimental data on the formation of order in a quiescent nematic polymer solution. It explains the concentration dependence of the nematic-isotropic transition temperature and deals with the effects of chain flexibility and molecular weight distribution. Moreover, the structure and properties of a nematic solution under shear and elongational flow have been clarified to a great extent.

6.3 MORPHOLOGY OF FIBERS AND FILMS

The structure and morphology of fibers and films can be described at various levels of structural dimensions, depending on the resolution of the diffraction equipment and of the microscopes used in the investigation. First the chain conformation in the solid state is discussed; then the packing modes of the chains in the crystallites determined by X-ray diffraction are reviewed; next the structural characteristics of the fibrils observed by X-ray and electron diffraction are examined; and finally the morphological features as seen by electron and optical microscopy are dealt with. After this survey of the structure and morphology, the formation of the fiber and film by coagulation from a lyotropic solution is discussed.

Fibers are identified here by their chemical names. In this respect it is useful to know that Twaron and Kevlar are the trade names for PpPT A fibers, and that PRD-49 is a fiber made from PpBA. The para-aromatic polyamide fibers PpPT A, PpBA T and PpBA are often called 'aramid' fibers.

6.3.1 Chain Conformation The relative importance of the various factors determining the persistence length in solution and the conformation in the solid state of aromatic polyamides are well illustrated by the differences between the para and the meta forms of poly(phenylene terephthalamide). Both chains consist of the same planar elements, viz. the phenyl and amide groups. The latter group adopts only the trans conformation in these polymers, wherein the chance of a transition to the cis conformation is extremely small owing to a barrier of about 60-80 kJ mol- 1. Although in both para- and metalinked chains the same intermolecular interactions exist between phenyl


and amide segments, the way in which these segments are joined is of primary importance for the conformation of the chain in solution. Thus, in the para-linked aromatic polyamides the directions of the rotation axes C(ar)-CO and C(ar)-NH are almost parallel for a number of successive units. (In crystals of low-molecular-weight amide compounds the valence angles at the carbonyl and the nitrogen are close to 116° and 1220

,

respectively.) Hence, irrespective of the degree of hindrance to rotation, the length of the chain increases progressively with the degree of polymerization, as a result of which appreciable values (20 nm) for the persistence length are obtained. This is not the case in the meta conformation where, as has been pointed out by Tsvetkov, the directions of rotating bonds can change by 60° in going from one monomer unit to the other. 98 This results in a persistence length of about 2·5 nm, which is typical of flexible-chain polymers.99

Apart from the inherently extended nature of the chain conformation of PpPT A, PpBA and PpBA T, other aspects of the conformation in the solid state are determined by competitive intramolecular interactions between the phenyl and amide segments. These are the resonance effect, tending to stabilize the coplanarity of these segments, and the counteracting steric repulsion between the oxygen and an ortho-hydrogen and between the amide hydrogen and an ortho-hydrogen. This results for PpPT A in an angle of - 30° between the amide plane and the terephthalic segment and an angle of 38° between the amide plane and the phenylene plane of the p-phenylenediamine segment. 100,101 These values for the internal rotation angles are close to those observed for the corresponding angles in benzamide/0 2 acetanilide,103 terephthalamide,104 N,N'-( p-phenylene)dibenzamide, 1 05 N,N'-diphenylterephthalamide/0 6 as well as in PmPTA.107,108 In the 4: 1 sulfuric acid complex of N,N'-(p-phenylene)dibenzamide the three phenyl rings are coplanar, but the amide groups are rotated out of this plane with torsion angles of 3r and _42°.109

Investigation of the conformational energies may provide more information on these interactions. However, different approaches do not yet give the same results: Laupretre and Monnerie found for the torsion angle OJ of the NH-CO group around the C(ar)-N bond in acetanilide a minimum of the conformational energy for OJ = 0°, indicating that steric repulsions are not evident in these calculations. ll0 ,111 By introducing the delocalization energy, Ed = - B cos 20J with B = 29 kJ mol-I, as well as steric repulsions and London dispersion attractions between nonbonded atoms, Hummel and Flory, however, calculated an equilibrium


angle of 30° with a rather low value for the barrier height of about II kJmol- l.112 Using the same contributions to the molecular energy, Tashiro et al. also arrived at an angle of about 30°, but with a barrier height of 25-50kJmol- l.113 This height seems reasonable compared with values estimated by NMR measurements for various model compounds.

In the PBT chain there is steric hindrance between a sulfur atom and a phenyl hydrogen atom, as well as between a nitrogen atom and a phenyl hydrogen atom. Conformational energy calculations by Welsh et al. yielded a rotation angle of 55° between the phenylene ring and the benzobisthiazole moiety.30.114 Taking into account also the intermolecular interactions, these authors arrived at a rotation angle in the range of 0-25°, which agrees with the value of 23° observed in a model compound.28 Calculations by Odell et al. gave a barrier height of 10 kJ mol- 1 in the planar conformation, which should normally ensure a non-planar conformation. However, a characteristic change in color from yellow to blue has been observed after heat treatment of films, which has often been associated with an increase in the degree of conjugation and thus of the planarity of the moleculeYs

The conformational energy calculations for the PBO chain by Welsh et al. resulted in a planar molecule, which appears to be consistent with experimental results. 30

6.3.2 Results of X-ray and Electron Diffraction Studies When a PpPT A fiber is obtained by air-gap spinning from a highly concentrated (20%) solution in sulfuric acid, the crystal structure formed after coagulation has a monoclinic unit cell (modification I) with pseudo-orthorhombic symmetry and two chains per cell (a =0,773-0,784 nm; b= 0·515-0·522 nm; c (repeat distance) = 1,28-1·29nm; y=90°, space group P2dn; crystalline density Pc= 1518-1542kgm- 3 ).100,10l,ll6 This structure is depicted in Fig. 6.11. Along the direction of the b-axis the chains are laterally bonded by hydrogen bonds, while weaker interactions exist in the other directions normal to the chain axis. Spinning from sulfuric acid solutions with a concentration less than 9% yields another crystal modification (II) identified by Haraguchi et al. ll 7,118 It has almost the same unit-cell dimensions but the hydrogen-bonded plane through the center of the unit cell is shifted along the b-axis over a distance of b/2. Spinning from solutions of 9-15% concentration yields fibers in which both crystal modifications coexist. 1 19


T

PpPTA PPBA , .. , ..

T - r-----, - ;

._+--"d'....,d . ~,Y /"

-t

~ ,-_::~r.=-.~~~ _. _ : ~ ~ .... _-: _

_ ~ b

(a) (b)

Fig. 6.11. Crystal structures of PpPTA (left) and PpBA (right).100.120

PpBA crystallizes in an orthorhombic unit cell (a=O·771 nm; b=O'514 nm; c (repeat distance) = 1·28 nm; space group P2 j 2 j 2b

Pc = 1540 kg m - 3). There are two monomeric units in the repeat distance and the phenyl groups are coplanar, subtending an angle of 40° with the amide groups. The polar chains are laterally bonded by hydrogen bonds along the b-axis (see Fig. 6.11). The unit cell


contains two chains, one up and the other downYo The length of the repeat unit of PpBAT is about 1·5 times that of

PpPT A, but the X-ray diffraction patterns are very similar, yielding the same unit-cell dimensions (a =0·797 nm; b =0·519 nm; c= 1·29 nm). Owing to the polarity of the 4,4' -diaminobenzanilide molecule, its amide group can take two orientations in the chain, leading to disorder in the crystal structure, which is the cause of the apparent repeat unit of 1·29 nm. The disorder also affects the lateral packing. Not more than five out of six possible hydrogen bonds per two monomeric units can be formed between the chains, which has been confirmed by IR spectroscopy.121

From the width of the reflections of the wide-angle X-ray scattering (W AXS) fiber pattern the apparent size and the degree of perfection of the crystalline domains can be determined. In PpPTA the size and disorder parameters are strongly anisotropic. Normal to the chain axis the dimensions of the domains are 4-6 nm. 122 Analysis of the meridional reflections shows that these fibers apparently have a paracrystalline structure with a longitudinal crystallite size of at least 25 nm and a lattice distortion parameter gIl of 2_3%.123 Annealing at temperatures up to 600°C increases the lateral size to about 10 nm and the longitudinal size to about 80nm, while gIl decreases to about 1%.124-126

An interesting correlation between the lattice distortion parameter and the modulus of PpPT A fibers was observed by Barton. 127 Extrapolation of the data to gIl = 0% leads to a modulus of about 218 GN m - 2, which is close to the value of the theoretical modulus of the PpPT A chain (see Section 6.4.1). Because the fiber compliance is linearly dependent on the second moment of the orientation distribution of the chains (see Section 6.4.1), Barton's observation indicates that the distortion parameter is probably related to the orientation distribution, i.e. an increasing disorientation of the chains will apparently result in a larger distortion parameter. This is supported by the relation observed by Hindeleh and Abdo between the modulus and a crystallinity index. By extrapolation they found a modulus of about 200 GN m - 2 at 100% crystallinity.128.129

Krenzer and Ruland applied a Fourier analysis to the meridional reflections of PpPT A. 130 They concluded that an interpretation of the line broadening by lattice defects in terms of a one-dimensional paracrystalline disorder along the c-axis is not justified, but that the disorder is the result of local defects such as chain ends incorporated in the crystalline lattice.


In the W AXS patterns of aromatic polyamides, often at least ten layer lines can be seen, which points to a high degree of order along the fiber axis direction. The overall appearance of the patterns is remarkably similar to the theoretically predicted patterns from curvilinear crystals discussed by VainshteinYI

The small-angle X-ray scattering (SAXS) pattern of para-aromatic polyamide fibers is characterized by an intense but diffuse equatorial scattering. This pattern has been interpreted on the basis of elongated microvoids oriented parallel to the fiber axis having an irregular crosssection with a dimension of 2-10 nm, whereas the longitudinal dimension is at least of the order of 25 nm but may be longer because of the limited resolution of the camera used in these studies. 125 ,132 An interpretation on the basis of a very broad distribution of lateral void sizes is probably more appropriate, as shown in Fig. 6.12(a) by the SAXS equatorial intensity curve measured with a linear position-sensitive counter up to a resolution of about 200 nm. Figure 6.12(b) presents the meridional SAXS curve of a PpPTA fiber, also measured with a linear position-sensitive counter up to a resolution of 190 nm. It shows clearly the absence of a periodicity in the electron density along the fiber axis with a wavelength smaller than the resolution and confirms the single-phase nature of the microstructure.1 33 This is very different from the SAXS patterns of fibers made of semi-crystalline polymers, such as polyamide and polyester, which show two- or four-point diffraction spots originating from a regular arrangement of amorphous and crystalline domains along the direction of the fiber axis that is associated with chain folding.

The exemplary electron diffraction (ED) studies by Dobb, Johnson and Saville122,134 deserve particular attention. Because ED patterns are taken from only a small volume of the fiber, they provide information about the local order and orientation distribution of the chains in the fiber. The ED patterns of PpPT A and PpBA show that within a certain domain size of about 0'5,um the crystallites are almost perfectly parallel-oriented. The meridional reflections have associated streaks along the layer lines, together with weak interference rings, indicating the presence of twodimensional lattices, although a large proportion of the material is in three-dimensional register. These two-dimensional lattices are formed by the hydrogen-bonded sheets, which have stacking faults along the chain direction.

It appears that the imaging of crystal lattice fringes is instructive in the visualization of the microstructure of the fiber. This was accomplished for the first time for an organic polymer with PpPT A


6.0

equator

5.0

• 4.0 • ::: •• .2 30

....

2.0

1.0

••• 0 -2.6 -2.2 -l8 -l4 -lO -0.6 -0.2

log( h)

(a)

6.0

meridian

5.0

4.0

~30

2.0

1.0

O+----.-----r----,----.-----.--~ -2.6 -2.2 -l8 -l4 -to -0.6 -0.2

log(h)

(b)

Fig. 6.12. Small-angle X-ray scattering curves of a PpPTA (Twaron) fiber measured with a linear position-sensitive counter by 1. Aerts, Akzo Res. Lab.,

h = 2n/d (d in Angstroms). (a) Equatorial intensity; (b) meridional intensity.

fibers by Dobb et aU 22 ,135 The lattice fringes are presented in Fig. 6.13 and show a very high degree of parallel alignment of the chains inside the crystallites, but very little evidence of paracrystalline lattice distortion, apart from the occasional appearance of curved layer planes. This suggests that Barton's observation 12 7 of the relation between the


<3 -(a) (b)

(c) (d)

(e) (I) Fig. 6.13. Optical and electron-microscopic micrographs of PpPT A fibers showing various features of the morphology. (a) ED-pattern; (b) lattice image of the 110 planes; (c) TEM-dark field image of the pleats, (d) SEM-image of filament cross-section, diameter l2,um; (e) peel morphology showing pleats; (f) fracture

morphology. Arrow indicates fiber axisY 1

distortion measured by X-ray diffraction and the fiber modulus indicates that at least a considerable part of the broadening along the scattering vector of the meridional reflections is, in fact, caused by the azimuthal spread of crystallites having meridional reflections with associated layerline streaks. This explanation concurs with the conclusion of Krenzer and Ruland,130 which is that there is mainly disorder caused by local defects.

304 M.G. Northolt alld D.l. Sikkema

The perfection, or absence of disorder, seen in the lattice fringes derived from equatorial reflections of aramid fibers is not found to be much different from the perfection of the meridional fringes. This may also be concluded from the lattice images of the PBT fibers. 136 Apparently the broadening of the equator reflections in the X-ray diffraction pattern is mainly caused by the lateral crystallite size distribution.

Several unit cells have been proposed for the crystal structure of PBT: a primitive monoclinic cell (a=0'583 nm; b=0'354 nm; c (repeat distance) = 1·235 nm; }' = 96°; crystalline density Pc = 1690 kg m - 3) and a non-primitive cell with two chains per cell (a=0·710 nm; b=0'665 nm; c (repeat distance) = 1·235 nm; }' = 63°; crystalline density Pc = 1690 kg m - 3) have been reported by Roche et al.137 A non-primitive cell with different dimensions was also proposed by Odell et al. 115 (a = 1·196 nm; b=0'355nm; c (repeat distance)=1·235nm). For heat-treated and oriented samples these authors found a primitive cell with the same dimensions as observed by Roche et al.

The electron diffraction patterns of heat-treated PBT show a very high molecular orientation. Discrete reflections are only observed on the equator, and layer lines are seen up to 20 orders. The general appearance of the diffraction pattern of PBT fibers can be regarded as a composite of what has been predicted for isolated periodic cylinders and for parallel arrays of infinite cylinders.136.138 Small random translations of the PBT chains along their axes cause the disppearance of all non-equatorial reflections. The absence of hydrogen bonds enhances the translational freedom and is the cause of lack of lateral registry. Hence, the scattering on the non-zero layer lines arises from the continuous squared Fourier transform of the single molecule. Crystallites are elongated in shape, with their length being about 5 times larger than their width, which ranges from 2 to 8 nm. Heat treatment of the fibers increases the lateral size to about 12 nm and the longitudinal size to about 60 nm.139

The SAXS pattern of uniaxial extruded PBT films revealed a network of oriented fibrils with irregular cross-sections, very similar to the structure found in fibers of the aramids. 140 The diffuse equatorial scattering, also observed, for example, in cellulose fibers, is typical of fibers manufactured by wet-spinning and air-gap spinning processes.

The crystal structure of PBO has been determined. 141.142 Two sets of dimensions for the primitive unit cell have been reported (a=0'565 nm; b=0·358nm; c (repeat distance)=1'174nm; y=102'5°) and (a=O'5598 nm; b=0'3540 nm; c (repeat distance) = 1·205 nm). Although the chain packing is similar to that of PBT, the electron diffraction


patterns of PBO fiber clearly indicate a three-dimensional lattice, which is virtually absent in the PBT fiber. This is probably caused by the fact that the PBO chains are planar compared with the non-planar PBT chains, thus promoting chain registry. Lattice images show that the chains run straight and nearly parallel to the fiber axis for more than 500 nm. Lattice distortion parameters normal and parallel to the fibre axis are about 1%. In contrast to other rigid-chain polymeric fibers, the transverse dimensions of the crystallites are greater than the axial dimensions.142.143

6.3.3 Characterization by Optical and Electron Microscopy A distinctive structural feature of PpPT A and other para-aromatic polyamide fibers is the pleated sheet, which is shown in Fig. 6.I3(c), (d) and (e). It was observed for the first time by Ballou in dark-field images taken from meridional reflections of the ED pattern. 144 This phenomenon has been studied extensively by Dobb et al. 134 The pleated sheet consists of parallel-oriented chains, which are hydrogen bonded along the direction of the crystallographic b-axis. They are usually oriented perpendicular to the surface of the filament, i.e. in the circular cross-section the b-axes are directed preferentially along the radius, as shown in Fig. 6.14. The angle between adjacent planes of the pleat is about 1700 and the distance between two pleats is about 250 nm, but both values may vary along the radius of the fiber cross-section.

The pleated sheet can also be observed with optical polarization microscopy.126.145 Figure 6.15 demonstrates the origin of the various band patterns in the polarizing microscope, which are seen on a longitudinal section of a fiber. In the dark areas the chains are aligned parallel to the analyzer or polarizer direction. Such a banded texture is often observed in films and fibers of lyotropic and thermotropic polymers.

The lateral texture can be studied by measuring the lateral birefringence in a thin fiber cross section,144.146 or by registration of the retardation profile across the diameter of a whole filament with an optical interference microscope. 147 149 The combination of high inherent polarizability and parallel orientation of the chains in the para-aromatic polyamides generates an unusually high birefringence (~n ~ 0·5). Calculated values for the refractive index ellipsoid of PpPT A are na = 1'51, nb= 1'73, nc =2'04, implying for the birefringence ~n=0-4I5 and for the lateral birefringence ~nl = 0,22.148 Hamza found slightly different values, nc = 2·199 and ~n = 0-457. 150 Table 6.2 lists the observed refractive indices of various aromatic polyamides.

306 M.G. Northolt alld D.J. Sikkema

Fig. 6.14. Radial arrangement of pleated sheets in a PpPT A filament. (From Ref. 134.)

Hagege et al. studied lateral texture in PpPT A filaments using the H2S-AgN03 penetration technique. 1s3 The Ag2S-impregnated cross-sections observed in optical and electron micrographs (TEM) show radial bands, which correspond to a regular radial stacking of areas of high permeability and areas of low permeability to H 2S and AgN03.

There is clear evidence of a skin-core structural differentiation in PpPTA fibers,126.143 which for a product of a wet-spinning process is not surprising. Figure 6.13(d) shows, for example, the change of the pleat spacing along the radius of the fiber cross-section.

Several models have been proposed for the structure of paraaromatic polyamide fibers. The simplest model, postulated by Northolt et al., is based on observations by X-ray and electron diffraction and is depicted in Fig. 6.16(a).12S.154.155 The fibre is considered as being built up of a parallel array of identical fibrils. Each fibril consists of a series of crystallites arranged end to end and the polymer chains run through these crystallites parallel to their symmetry axes, which


+10

o

- 10

a

b

c

P

).c"A " \

d

Fig. 6.15. Banded textures under the polarizing microscope and the origin for the banded texture in a longitudinal section of PpPTA fiber under crossed polars (P and A). Chain direction in domain A at + 10°, domain B parallel to and domain C at _10° with respect to the fiber axis. Optical micrographs of banded textures are shown for orthogonal positions of the polars at + 10°, 0°, and _10°

deviation from the fiber axis. 151


Table 6.2 Observed refractive indices of para-aromatic polyamide fibers and a PBO fiber; nr radially lateral; nt tangentially lateral; n~ lateral; nil longitudinal. The results of Weeda were obtained by matching immersion liquids under the interference microscope, or from measuring the retardation. lsi For PBO fiber, nil was

estimated from a birefringence determination

Polymer (Ref) nr n~

PpPT A (as spun) (124) 0-474 PpPTA (600°C) (124) 0·508 PpPT A (Kevlar) (150) 1·626 2·275 0·649

(149) 0·026 0-435--0·461 (145) 0·45

PpPTA (Kevlar 49) (150) 1·605 2·267 0·662 PpPTA (Kevlar 149) (151) 1·681 1·598 2·13 PpPT A (Twaron) (151 ) 1·671 1·583 2·06 0·083 PpPTA (Twaron HM) (151) 1·679 1·598 2·08 0·087 PpBA (PRD49) (150) 1·644 2-405 0·761

(152) 0·65 PpBAT (151) 1·667 1·608 2·08 PBO (151) 1·663 1·589 >3-0

follow an orientation distribution with respect to the fiber axis. As demonstrated in Section 6.4.1 this model provides expressions for the compliance and the elastic tensile curve of the para-aromatic polyamide and other high-modulus fibers which have been experimentally confirmed. Although elongated voids exist between the fibrils, they are not completely separated and weak cross-links between the fibrils are formed by, for example, hydrogen bonds. Presumably disruption of this interfibrillar network causes the slight effect of a yield point in the tensile curve of the para-aromatic polyamide fibers.

The models of Morgan et al. 156 and Panar et al. 126 emphasize, among other things, the structural features as observed on etched specimens of PpPT A fibers, but do not provide a quantitative interpretation of the tensile properties of these fibers. Morgan postulates that the chain-end concentration and its distribution within the fiber are the principal structural factors affecting the deformation and failure processes and strength of PpPTA fibers. In the model depicted in Fig. 6.16(b) the chain ends in the skin are arranged essentially randomly relative to one another but become progressively more clustered in the fiber interior, resulting in periodic transverse weak planes about every 200 nm along the fiber axis, supposedly being the average chain length. The fiber core consists of


(a)

Core

a

I f t , \

a

\\-60nm Skin 0.1 -1 P m thick random cha in end di s1 ribu tion

1 Fibril ax is

(b)

Shish- Kebab

Fibril = == == _ Ordered =.: == lamellae ---- -------Del",,! zone = ::'== 1_ Ti~ _ __ • pomt

--------------,---------------------------------------------... .-. --. --_ . - "":.1;;; 1--nm

(el

Fibre ax is

Fi ll,., oxi!3

'-Chain with I.mgth 220 nm

Fig. 6.16. Models for the structure of aromatic polyamide fibers: (a) Northolt; (b) Morgan; (c) Panar.

cylindrical crystallites about 60 nm in diameter. Panar's model, derived from investigations of the same kind of fiber, is presented in Fig. 6.l6{c). It shows fibrils with a diameter of 600 nm (ten times larger than Morgan's fibrils) with regularly spaced defect zones oriented normal to the fiber


axis with a period of 30-40 nm, presumably corresponding to an average chain length. In both models the pleated sheet, as observed by Dobb et al., is thought to be superimposed on the fibrillar structure. Panar suggested that the pleats impose sufficient strain on the chains that preferential chemical degradation at these locations is responsible for the visualization of the defect planes after an etching treatment. Li et al. made secondary-electron images of ion-etched PpPT A and PpBA fibers; they also found a layered structure with a 20-30 nm spacing perpendicular to the fi ber axis. 1 57

Finally, the model of Termonia and Smith should be mentioned here, although it is not directly based on results of diffraction and microscopic studies.158160 It deals with the strength of high-modulus fibers, not with the elastic properties, and is discussed in Section 6.4.3.

Visualization of PBT films and fibers by electron micrography shows microfibrils of about 10 nm width. 161 As-spun fibers contain a large number of voids of elongated shape, and scanning and transmission electron microscopy showed that fibrils and microfibrils are predominantly lath-shaped. 162 In dark-field images a very marked banding transverse to the fiber axis has been observed, which probably originates from the pleated sheets seen also in fibers of aromatic polyamides. The periodicity of the banding is about 120 nm. 13 7

Young et al. performed a detailed study of PBO fibres using electron microscopy and Raman microscopy.143 Micrographs of longitudinal sections taken through as-spun and heat-treated PBO fibers show a fibrillar structure on a 100 nm to 200 nm scale. Unlike the aramids, no banding pattern has been observed. Skin--core differences similar to those formed in aramid fibers were also found in the PBO fibers. As-spun and heat-treated fibers showed a better orientation in the skin than in the core.

We conclude that there is a great similarity in structure at all levels between the different polymer fibers spun from lyotropic solutions. Moreover, many features of this structure have also been found in well-oriented cellulose fibers. 163, 164

6.3.4 Coagulation and Structure Formation The actual formation of the fibers and films in a spinning or extrusion process takes place during coagulation of the lyotropic solution. In this process step a phase transition to the solid state is induced by a non-solvent. The conditions of the coagulation process, e.g. the concentration of the polymer solution, the composition and the temperature of


the coagulation medium, determine to a large extent the morphology, texture and crystal modification of fiber and film. Complications arise because the spinning dope solidifies just before the coagulation starts. Hence coagulation is probably the least understood stage in the spinning process of lyotropic polymers.

In order to understand the coagulation process, the phase diagram of the polymer - solvent - non-solvent system and the solvate structure, being the precursor of the fiber structure, should be known. In a review paper by Iovleva and Papkov the results of investigations into polymersolvate crystals, in particular of aromatic polyamides, are generalized and phase diagrams are proposed. 165 Irrespective of the interactions whereby polymer-solvate crystals are formed (hydrogen bonds, interaction between the acid and the amide groups with protonation of the nitrogen), they are relatively unstable at elevated temperatures and the melting temperatures therefore lie in the range 50-130°C. The melting point depends on the ratio of the polymer to solvent, e.g. for a PpPT A -sulfuric acid solution it drops from about 80G C at 20% (w/w) to about 40°C at 10% (w/w) PpPTA. Papkov et al. observed in the PpBA/sulfuric acid system that the polymer-solvate crystals having spherulitic morphology coexist with a nematic and an isotropic liquid phase,166 which illustrates the complex character of the phase diagrams of solutions of lyotropic polymers. In the phase diagram for low-molecular-weight PpPT A in sulfuric acid, three regions have been found. At low concentrations the stable structure is that of an isotropic liquid. At slightly higher concentrations the solutions are clearly biphasic, with coexisting isotropic and nematic domains. At still higher concentrations the system becomes homogeneously nematic at high temperatures and forms an ordered solvate crystal at low temperatures. 167

Pertinent data on the solvate crystals of PpPTA with sulfuric acid have been given by Arpin et al. 168 They found two crystal modifications depending on the thermal treatment applied to a nematic solution of PpPTA in sulfuric acid. The lattice spacings observed by Iovleva et al. 165

and Xu et al. 169 for these solvate crystals are identified as mixtures of the two forms found by Arpin et al. It is remarkable that the X-ray diffraction patterns of these polymer-solvate crystals show lines that are in general much sharper than those of the pure polymer, indicating that the solvate crystals are composed of large and rather perfect crystallites. This has also been observed for the X-ray diffraction patterns of filaments prepared by spinning a nematic solution of PpPT A in sulfuric acid into a non-coagulating bath.


Unit-cell dimensions of the PpPTA- and PpBA-solvate crystals with sulfuric acid have been determined by Gardner/67.170 but do not yield the lattice spacings observed by Arpin et al. The structures proposed by Gardner are made up of hydrogen-bonded sheets consisting of polymer chains alternating with sulfuric acid molecules. In the PpBA solvate the sheets are closely packed, whereas the PpPT A solvate has additional acid molecules separating the sheets. The crystal structure determination of the 4: 1 sulfuric acid - N,N'-( p-phenylene)dibenzamide complex supports these solvate models and shows that the amide groups are protonated. 109

Detailed investigations of the coagulation of PpPT A films from sulfuric acid solutions in various coagulants have been carried out by Haraguchi et al.117.118 They found that films coagulated in organic liquids showed the crystal modification I with a uniplanar orientation and the hydrogen bonds directed parallel (P) to the film surface. 1 00 On the other hand, films coagulated in water had modification II, with a uniplanar orientation in which the hydrogen bonds are oriented normal (N) to the film surface. The two kinds of films showed different mechanical properties (see Section 6.4.3). The formation of these two kinds of films is explained by the anisotropy in the tendency to aggregation of the PpPT A chains due to the difference in interactions between polymer-polymer and polymercoagulant during the coagulation process. Upon annealing, only the film with modification II and N-texture will undergo a crystal transformation, namely into modification I with P-texture. This only happens for films coagulated in water and prepared from solutions of low polymer concentration. All these films were prepared by smearing on a glass slide and soaking in a coagulant. Unpublished results from our laboratory showed that only in films prepared from anisotropic solutions is the orientation of the PpPT A chains parallel to the shear direction. Negative PpPT A spherulites were prepared by Takahashi et al. from a 15% sulfuric acid solution.! 71 Shearing of the solution produced a banded texture in the coagulation product, with bands of 43-77 nm width that are oriented normal to the shear direction with the preferential direction of the chains parallel to this direction (see also Section 6.2.2). Similar phenomena were observed when PpPT A was coagulated from an organic solvent. 1 72 Xu et al. studied PpPT A spherulites prepared from solvate spherulites that were grown in a 12% sulfuric acid solution at room temperature and washed afterwards. These were positive spherulites with the modification II crystal structure. 173

In a recent study Roche et al. performed coagulation experiments with films prepared from a 20% (w/w) solution of PpPTA in sulfuric acid. 174


They also found that the banding pattern arises only after cessation of flow and that its periodicity depends on the shear rate and film thickness. Row-nucleated morphologies were obtained when quenching was performed in a non-coagulating liquid, allowing the formation of a crystallosolvate as an intermediate state. The lamellar thickness of these structures observed by small-angle electron diffraction ranged from 40 to 240 nm and was found to be linearly related to the average chain length.

With regard to fiber formation, these authors conclude that relaxation of the spinning dope before coagulation is the cause of the banding pattern, its periodicity being determined by the coagulation rate. Coagulation conditions similar to fiber spinning never showed the row-nucleated morphology. However, only slow diffusion of the coagulant yielded the modification II structure, which is presumably formed after crystallization and subsequent coagulation of the crystallosolvate.

As has been mentioned, PpPT A fibers spun with an air gap from a sulfuric acid solution into a water bath show lateral texture, i.e. the hydrogen bonds in the fiber cross-section are oriented perpendicular to the circumference of the filament cross-section. As shown by Van der Zwaag, the degree of lateral texture in as-spun PpPT A fibers can be continuously varied between random and radial texture by changing the composition of the coagulation bath. 17 5 Using the optical interference microscope and a range of immersion fluids, the refractive index in the radial, nn and the tangential direction, nt, can be measured on whole filaments and the difference l1n ~ = nr - nt, called the lateral birefringence, is a measure of the degree of lateral texture. Figure 6.17 shows that the

005

0.04

003

-l c '1

0.02

0.01

0 0 5 10

• ... • • .

15 20 25

Solubility parameter (eal lem3 )'"

Fig .. 6.17. Lateral birefringence of the fiber versus the solubility of the coagulation liquid. 175


solubility parameter of the coagulation liquid is strongly correlated with the degree of lateral texture of PpPT A filaments, which varies from ~n1 =0 for random texture in the cross section to ~n1 =0·045 for a well-developed radial texture.

Spinning of PpPT A/sulfuric acid solutions in water showed that for a spinning dope with a low polymer concentration (3-9%) only modification II is formed in the fiber. For medium concentrations (10--15%) both phases coexist in the fiber, while for high polymer concentrations (> 15%) only modification I is formed. 119 This is different from the results obtained with the coagulation of films in water, whereby modification II is always formed. Air-gap spinning of a 19·5% PpPTA solution in water, with and without a drawing stress, showed that the external stress suppresses the formation of modification II during coagulation. Raising the acid concentration of the coagulation bath increases the amount of modification II in the fiber at the expense of modification I. The opposite was observed for the coagulation of films in water. Increase of the acid concentration in the coagulation bath also raises the void content in the fiber as observed by SAXS and TEM.1 76

As discussed previously, sheared lyotropic solutions exhibit banded textures, while very similar textures are also observed in films and fibers after coagulation. Chen et at. concluded that coagulation from isotropic solutions yielded only spherulites, whereas spherulites, shish-kebabs and banded textures can be grown in films coagulated from sheared anisotropic solutions. 177 The width of the bands in the films appeared to depend on the speed of coagulation, varying from about 600 nm for fast coagulation to about 8 J.lm for very slow coagulation. Picken found that banded textures in nematic solutions of aromatic poly ami des are only formed immediately after cessation of the shear flow. These banded textures are frozen in by coagulation. When relaxation is allowed by slow coagulation, no bands are observed. 178 This has also been observed in mesomorphic hydroxypropyl cellulose solutions. 1 79

Horio et al. found that the band spacing in fibers spun from a 20 wt% PpPT A solution changed from 300 to 530 nm as the sulfuric acid concentration in the water bath increases from 0 to 60%.180 Washing and drying did not have any effect on the patterns. The authors suggest that elastic rebound due to the velocity gradient of the fiber in the bath and the coagulation rate determine the spacing of the banded texture. Presumably the increase of the pleat spacing from the skin to the core in the filament, as shown in Fig. 6.13, may also be caused by a recoil effect due to drag of the coagulation liquid on the spinline and by a difference


in coagulation rate resulting from the acid concentration gradient along the filament radius.

A study on the coagulation of PpBA films from solutions in N,N'dimethylacetamide with Liel in a magnetic field was reported by Takase et al. 18 1.182 Good uniaxial orientation with the chains aligned parallel to the field was obtained with fields exceeding 1 Tesla or IT. Two crystal modifications of PpBA were found, depending on whether the crystal solvate was annealed before or after loss of the solvent.

In the spinning process the aramid fibers are thoroughly washed after coagulation and subsequently dried. Application of tension during coagulation affects the mechanical properties, in particular it increases the modulus of the fiber. This is also true for the drying stage in the process and for an optional heat-treatment of the fibers.

The coagulation bath employed in the spinning of PBT fibers and films is typically a mixture of water and methane sulfonic acid when the spinning dope consists of a solution of methane sulfonic acid (MSA), and it is water for spinning solutions of poly phosphoric acid (PPA). As has been observed with spinning of aromatic polyamides fibers, lower coagulation rates and/or lower bath temperatures have produced fibers with higher modulus and strength. The mechanical properties of the fibers spun from MSA dopes were inferior to those obtained with PPA dopes, owing to large voids.31.183.184

6.4 MECHANICAL AND THERMAL PROPERTIES

The most prominent physical characteristics of the lyotropic polymers are the tensile and thermal properties of fibers made from these polymers. Fibers of para-aromatic polyamides, PBO and PBT are characterized by high tensile strength (crb>2GPa), low elongation at break (sb<6%) and high-to-ultrahigh modulus (50 < E < 400 GPa). A survey of these properties is listed in Table 6.3.

Fibers and films possess anisotropic mechanical properties. Hence a discussion of this subject should include tensile and compression properties in at least two different directions, viz. in longitudinal and transverse direction to the filament axis. However, in general little is known of the transverse properties of fibers and films.

The discussion of mechanical properties includes the various contributions of elastic, viscoelastic and plastic deformation processes. Often two characteristic stress levels can be defined in the tensile curve of polymer


Table 6.3 Modulus (E), strength (O"b)' elongation at break (eb) of filaments and the density

(p) of para-aromatic polyamide, PBO and PBT fibers

Fiber type p E O"b eb Test length (Ref) a

(kgrn- 3 ) (GNrn- 2) (GNrn- 2) (%) (rnrn)

PpPTA Twaron 1442 91 3-47 3-4 250 PpPTA Twaron 1454 120 JlO 2·1 250

HM PpPTN 1440 67 J50 4·7 100 PpPTAb 1440 140 4·09 JO 100 PpBAT as-spun 1423 102 2·87 J2 100 PpBAT heat-tr. 1433 152 2·75 2·0 100 PBO as-spun 1500 80 2-31 2·1 250 PBO as-spun 144 4·6 J2 50 143 PGO 600°C 250 5·1 1·9 50 143 PBO 650°C 262 3-4 1-3 50 143 PBT 1470-1530 18 2·35 7·1 25 31 PBT 66 2·28 4·8 25 31 PBT 159 2·35 2A 25 31 PBT 1540-1600 303 3-49 1·3 25 31 PBT 331 4·19 1-4 25 31

alf no reference is given, results are from Akzo Research Lab., strain rate 5% min-I. bExperimental fibers.

fibers: the yield stress, at which a significant drop in slope of the stress-strain curve occurs, and the stress at fracture, usually called the tensile strength or tenacity. In this section the relation is discussed between the morphology of fibers and films made from lyotropic polymers and their mechanical properties such as modulus, tensile strength, creep and stress relaxation.

Other properties such as fatigue, knot strength and abrasion resistance are based on a combination of the various anisotropic mechanical and thermal properties. Their relation to the fiber morphology is rather complicated and is still not well understood.

The tensile curves of a series of PpPT A fibers with different initial moduli are shown in Fig. 6.18. Two ranges can be distinguished, viz. a short and straight part up to a kind of yield point at a strain of about 0'5%, and an extended concave curve up to fracture constituting the major part of the tensile curve. Presumably up to the yield strain the fibrillar structure is extended without disrupting the interfibril hydrogen bonds. 164 Approaching the yield point, rupture of these bonds takes place


4~---------------------'

~3 z ~ 00 2 '" l' iii

O~--~--~----~--~--~

o 3 Strain(%)

317

Fig. 6.18. Tensile curves of PpPT A mono filaments with different initial moduli, drawing speed 10% min - I, test length 10 cm.

and a different deformation process is initiated. This is accentuated by the difference of the various tensile curves in the slope before and after the yield point, which is clearly a strain-determined phenomenon and similar to that observed in cellulose fibers.163 The transition is easily observable for the low-modulus fibers in this figure and disappears with increasing modulus. In the second stage of the deformation process a steady increase of the dynamic modulus is observed which, as will be shown later, indicates a progressive contraction of the orientation distribution of the chains. 154,185 This process may also be interpreted as a stretching or 'straightening' of the fibrils.

An important aspect of the tensile deformation of fibers of semicrystalline and crystalline polymers is the recovery after unloading. When the first extension of an oriented fiber reaches well into the second stage of the tensile curve, i.e. beyond the yield point, then the recovery is not complete. The permanent extension or set is approximately equal to the extension at which yielding occurs. The recoverable extension shows a spontaneous and latent recovery corresponding to elastic and viscoelastic or delayed elastic contributions. Further repeated extension of the fiber up to the same maximum extension hardly increases the permanent deformation but is still accompanied by a little hysteresis, as shown in Fig. 6.19.

The tensile curves of PBT fibers presented in the literature show a stress at rupture of only about 1 GN m - 2, whereas the reported tenacity values (see Table 6.3) are much higher. 162,186 This implies that the tensile curves shown in the literature do not extend beyond the yield point near 0·7%. Figure 6.20 shows the tensile curves of first and multiple loadings of an experimental PBO fiber. The yield point near 0·8% is noticeable. The great similarity between the tensile curves of the aramid and PBO


3~---------------------------------------'

E 2 z <!l rJl rJl

~ 001

2/0 1 Strain(%)

2 3

Fig. 6.19. Tensile curves of the first (solid curve) and the third (dotted curve) cycle of a standard and high-modulus PpPT A monofilament, drawing speed

5 % min - 1, test length 25 em.

fibers (and presumably PBT fiber) indicates the same deformation process, which with regard to the elastic part is described in Section 6.4.1. As with the aramid fibers, the Young's modulus increases after first loading only, owing to a slight permanent contraction of the chain orientation distribution. The effect seems to be much more noticeable in the tensile curves of the PBT fiber, perhaps owing to a larger void content of these fibers. The empty space of the voids offers ample room for rotation of fibril segments towards the stress direction. Therefore, the permanent contraction of the orientation distribution after first loading should be much smaller after heat treatment of these fibers under tension.

Also typical of fibers made from lyotropic polymers is the fact that the hysteresis observed during cyclic loading is very small compared with that of fibers of semi-crystalline polymers. The dissipated energy relative to the stored energy is 7% for PpPT A fibers with moduli up to 140 GN m - 2 and 12% for the experimental PBO fiber with a modulus of 150 GN m - 2, while for a well-oriented poly(ethylene terephthalate) fiber (tire yarn) with a modulus of 18GNm- 2 this ratio is 47%.

6.4.1 Elastic Behavior Various theories of the elastic extension of oriented fibers have been developed. In a series model, uniform stress throughout the fiber is assumed. Two kinds of series models have been proposed. The classical series model, first applied by Ward 187 to polymer fibers, is based on a series arrangement of cube-shaped elements, whereas in the modified


2.5,-----------------,

2.0

~15 C!J

'" ~ 1.0

iii

0.5

O~-_.--._--r_-_r-~

E Z

o 0.5 to 1.5 2.0 2.5 Strain ('!o)

(a)

1.6,..---------------,

1.2

C!J 0.8

'" '" ~ iii

0.4

O~-L---.----.---~

o 0.5 to 1.5 Strain ('!o)

(b)

319

Fig. 6.20. Tensile curves of an experimental as-spun PBO monofilament, drawing speed 5% min -1, test length 25 cm. (a) Curve up to rupture; (b) first (solid

curve) and third (dotted curve) cycle.

series model proposed by Northolt and van der Hout oblong-shaped elements are used. 155 The modified version of the series model, which takes account of the anisotropic and fibrillar structure, has been further developed to give the elastic tensile curve of oriented and crystalline polymeric fiber. 15 5

Recently Allen and Roche l85 presented equations for the modulus and the tensile curve of a fiber that had already been derived earlier by Northolt. 154 However, in both papers the distinction with respect to the


shape of the elements pertaining to the microstructure has been overlooked.

Other possible models for fibers are the uniform strain and the elastic unwrinkling model. The first does not describe the observed linear relation between the dynamic compliance and the orientation distribution parameter of the chains in the polymer fibers; the second is not essentially different from the classical series modeU 88

Here, only the modified series model by Northolt and van der Hout is discussed. According to this model the fiber is regarded as being built up of a parallel array of identical fibrils which are subjected to a uniform stress along the fiber axis. Each fibril consists of a serial chain of crystallites. A crystallite is composed of rigid-rod polymer chains running parallel to the symmetry axis (see Fig. 6.l6(a). All crystallites are considered to be transversely isotropic and have identical mechanical properties. The orientation angles, <p, of the symmetry axes of the crystallites relative to the fibril axis follow a distribution f(<p) along the meridian. The elastic extension of the fibril is the result of the distortion of the crystallites. The latter is determined by two mechanisms, viz. the extension of the chain and the shear between adjacent chains. The contribution of each individual crystallite to the fibril strain is a function of its orientation angle q;. For this modified series model of a fiber it can be shown that the compliance, S, is given by

I <sin2 q;) S = - + --,-----:-----'--'-

ec 290 (6.15)

where ec is the chain modulus, 90 the modulus for shear between adjacent chains and <sin2 q;) the second moment of the distribution defined by

f/2 f(q;) sin 3 <p cos q; dq;

<sin2 <p) = fO nl2 (6.16) o f( q;) sin q; cos q; d<p

For the relatively narrow orientation distributions commonly found in polymer fibers the second moment defined according to eqn (6.16) hardly differs from the one calculated with the standard averaging procedure.

The equation for the elastic tensile curve of the fiber with well-oriented chains is

(6.17)


where <sin2 <Po) is the initial value of the orientation parameter before loading. Equation (6.15) for the compliance shows that the elastic response of the fiber is equal to that of a serial arrangement of two springs, viz. a 'chain spring' with modulus ee and a 'shear spring' with modulus 2go/<sin 2 <p). The chain spring provides the chain stretching contribution, Se = (Jlee, while the shear spring imparts the rotational contribution to the fiber strain.

<cos <p) - <cos <Po) Sr = --<-c-o-s -<P-o )--

This mechanical model is depicted in Fig. 6.21. Equations (6.15) and (6.17) have been confirmed for aromatic polyamide fibers by a variety of experiments. Figure 6.22 shows the dynamic compliance versus the orientation parameter measured during extension of medium and highmodulus PpPT A fibers. It confirms the linear relation (6.15) and yields ee = 240 GN m - 2 and go = 2 GN m - 2. Furthermore, the tensile curves of the second and higher extensions of an aramid fiber are well described by eqn (6.17).155 The theory also yields a relation between the strain and the dynamic modulus, which agrees well with the experimental results.163.164

A different test of the mechanical model is provided by measuring the resonance Raman scattering during fiber extension. Raman bands

a_

Compliance: S =..l.. + <sin2(JJ> ec 2 go

a <sin2p > [ J Strain: E : e; + 2 a 1 - exp(· alga)

E : f (chain) + f (shear)

Simple series model

Fig. 6.21. Schematic presentation of the mechanical model of an oriented and crystalline fiber: a chain spring in series coupling with a shear spring.


15r---------------------------------~

o+--------.------~._------._------~ o 0.01 0.03 0.04

Fig. 6.22. Dynamic compliance S measured at 10 kHz as a function of the orientation parameter during extension of different PpPTA fibers; open symbols

for the first extension and filled symbols for the second extension. 154

assQciated with vibratiQnal mQdes of atoms in the PpPT A chain are strain dependent, as shQwn by GaliQtis et al. 189 AccQrding to. eqn (6.17) the ratio. Qf the chain extension to. the rQtatiQnal strain contribution varies with the initial QrientatiQn parameter Qf the fiber, i.e. the cQntributiQn due to. chain extensiQn, Ce, increases with increasing fiber modulus. Since the shift in the Raman spectra is due to. the stretching Qf the cO. valent bQnds in the chain, the shift should be proportional to. Ee Qnly and independent of Er . FQr fQur different PpPT A fibers with mQduli ranging from 49 to. 135 GN m - 2 the shift in frequency Qf the 1610 cm- I

Raman band per percentage applied strain was measured. The increase of this shift with the fiber mQdulus was fQund to. be in agreement with the theQry.190 The same effect was also. Qbserved by Young et al. fQr PBO fibers. High-mQdulus fibers Qbtained by heat treatment shQwed a larger shift than the IQw-mQdulus as-spun fibers. 143 AccQrding to. YQung this phenQmenQn is generally Qbserved fQr fibers such as PBO, PBT, aramids and carbQn fibers, implying that the elastic behaviQr Qf all these fibers can be interpreted by the unifQrm stress mQdel.

As shQwn by eqns (6.15) and (6.17), the tensile elastic behaviQr Qf fibers made from lyQtrQpic PQlymers is determined by the chain mQdulus ee, the shear mQdulus go and the QrientatiQn parameter <sin2 ({Jo). As has been discussed in SectiQns 6.2 and 6.3, the latter is determined by the persistence length, the mQlecular weight, the PQlymer cQncentratiQn and


the temperature of the liquid-crystalline solution, and furthermore by the spinning and coagulation conditions.

We will now turn our attention to the determination of the intrinsic polymer properties, ec and go. The first theoretical estimate of the chain modulus of aromatic polyamides was carried out by FieldingRussell. 191 Using Treloar's method,192 he arrived at a value of 200 GN m - 2 for both PpPTA and PpBA, and a value of 127 GN m - 2 for PmPT A and PmBA. Tashiro et al. calculated for the chain modulus 182GNm- 2 for PpPTA, 163GNm- 2 for PpBA, and 90GNm- 2 for PmPTA. 113 Using molecular mechanics, Kooijman et al. found for the PpPTA chain a value of 204GNm- 2.193 Measuring the lattice strain along the fiber axis with X-ray diffraction, Gaymans et al. found an experimental value for the PpPT A chain modulus of 200 GN m - 2.194 As mentioned previously, Barton127 found 218 GN m- 2 from an extrapolation of the experimental relation between the lattice distortion and the fiber modulus, and 240 GN m - 2 was found by regression analysis from Fig. 6.22.154 Measurement of the lattice strain as a function of temperature yielded the relation between the chain modulus and temperature:195 for PpBA it was found that ec =

188GNm- 2 at room temperature and 235GNm- 2 at OK from extrapolation. For PpPTA these values are 168 and 200 GN m - 2, respectively. The general agreement between the calculated chain modulus and the observed lattice modulus along the chain direction (also found for other linearly extended polymers) provides substantial evidence that interchain cohesion has at the most only a minor effect on the lattice extensibility in the chain direction.

Kooijman et al. calculated the modulus for shear between adjacent chains in the hydrogen-bonded plane of PpPT A and found s 551 =

4·1 GN m - 2.193 This value agrees well with the calculated modulus for the shear between adjacent chains in the hydrogen-bonded plane of nylon 6.196 The monoclinic (pseudo-orthorhombic) PpPT A crystal has two moduli for the shear between adjacent chains, viz. s;;i and s5l. Because of the absence of strong secondary intersections, such as hydrogen bonds, between the chains in the 010 plane, S44 will be larger than S55. Hence, it is not surprising to find for go a value of 2GNm- 2, which agrees well with the observed shear or torsion modulus of PpPT A single filaments: G= 1·8 GN m-2.197-199

As discussed previously, para-aromatic polyamide fibers show a pleated sheet, which can also be described as a sinusoidal undulation of the fibrils. The effective orientation parameter determining the modulus


of the fiber may then be approximated by

(sin2 <iO)= (sin2 <iO p )+ (sin2 <iOf)

where (sin2 <iO p ) is determined by the distribution of the pleat angle and (sin2 <iOf) by the orientation distribution of the crystalline axes in the non-undulated fibrils. Observation of a PpPT A fiber during tensile deformation in the polarization microscope shows the disappearance of the undulation at higher strains. It has also been observed that the pleat angle increases for PpPTA fibers with increasing modulus.1s5 The pleated sheet is absent in fibers with a modulus of 179 GN m - 2.200

Determination of the crystal lattice strain along the c-axis by X-ray diffraction for a PBT fiber yielded a modulus of 395 GN m -2, and for PBO fiber 460 GN m - 2. Theoretical calculations of the chain moduli resulted in much larger values, viz. 615 and 730 GN m- 2, respectively.201

As-spun PBT fibers reach static moduli of about 190 GN m - 2 and a tenacity of 2·1 GN m -2.201 Heat-treatments at temperatures between 630 and 680°C with tensions in the range 0·1-0· 2 GN m - 2 produce fibers with moduli up to 335 GN m - 2 and tenacities up to 3·9 GN m - 2.201,202 The largest reported values for as-spun PBO fibers are 179 GN m - 2 and 4·6 GN m - 2, whereas a heat treatment yielded a modulus of 345 GN m - 2 and a tenacity of 5·8 GN m - 2.143,201 Aramid fibers differ in this respect, since only the modulus can be substantially increased by a heat treatment: the maximum tenacity values for both as-spun and heat-treated filaments, measured for a test length of 25 mm, are close to 5 GN m - 2.

6.4.2 Creep and Stress Relaxation The creep of polymer fibers is composed of the primary creep, which is recoverable with time, and of the non-recoverable or secondary creep.203-205 The secondary creep is almost negligible if a fiber has been mechanically pre-conditioned, i.e. when it has first been stretched to a strain larger than the strain that will be ultimately reached in the subsequent creep experiment. Thus the results of creep measurements can only be interpreted as true viscoelastic measurements when performed on mechanically pre-conditioned fibers.

Creep measurements on single PpPTA filaments have been performed by Walton and Majumdar.206 They observed creep strains amounting to less than 20% of the initial elastic strain after several years under stress. The graphs of the creep plotted against the logarithm of time were not linear but showed an increasing slope. Ericksen207 observed logarithmic creep curves and, in addition, found that the strain after recovery from creep also


followed a logarithmic time law. Furthermore, the creep rate after first loading was higher than the rate observed if the specimen had crept, was allowed to recover, and was then reloaded. This means that the creep rate of mechanically conditioned fibers is lower than that of unconditioned fibers.

Creep and stress relaxation rates of para-aromatic polyamide fibers are extremely low compared with those of conventional fibers like nylon and polyester. For example, the creep rate of an unconditioned PpPT A fiber increases from 2·5 x 10- 4 for 0·4 GN m - 2 to 7·5 X 10- 4 per decade for a stress of 2·8 GN m -2. Presumably this is caused by the combination of the rigid nature of the chains, the high crystallinity of the structure and the hydrogen bonding between the chains. By investigating the structural changes during creep and stress relaxation Northolt et al.208.209 arrived at a model for viscoelastic extension of well-oriented fibers. In these fibers the dynamic compliance, S, was found to be linearly related to the orientation parameter <sin 2 <p) according to eqn (6.15). Thus by measuring S during creep and stress relaxation the change of the orientation distribution can be monitored.

The modified series model shown in Fig. 6.21 has been extended to include the viscoelastic behavior. To this end the simple assumption is made that the time-dependent part of the creep strain arises solely from the rotation of the chains towards the direction of the fiber axis as a result of the shear deformation of the crystallites. This yields for the fiber extension as a function of the time t during creep caused by a stress 0"0:

( 6.18)

At any time the orientation parameter is given by the equation for the dynamic compliance (6.15) and we obtain

(6.19)

where So is the dynamic compliance before loading. A similar relation can be derived for the stress relaxation and we find, irrespective of the functional dependence on the time, the following relations:

*)-e(to) S(to) - S(t) go

(6.20)

O"(to)- O"(t) S(to) - S(t) = goec

(6.21)


These relations imply a progressive contraction of the orientation distribution in the fiber during creep and stress relaxation.

Figure 6.23 demonstrates the coupling that exists between the strain and the dynamic compliance during creep and recovery of a PpPT A fiber. During creep a gradual contraction of the orientation distribution of the chains in the fiber takes place. A similar result is shown in Fig. 6.24 for the stress relaxation. All these experiments have shown a logarithmic dependence on the time. Furthermore, it was found that for low and medium tensions eqns (6.20) and (6.21) are obeyed, but for high tensions and imposed strains the observed values of -As/AS and Au/AS are larger than the theoretical values. This discrepancy is probably caused by secondary creep, which becomes significant again as the creep stress approaches the stress used in the preceding mechanical conditioning procedure. Presumably its mechanism is a slip movement of adjacent chains during the orientation process and can be well visualized by a row of books that is slowly falling over.

1.25 8.0

~ ............. ~cco 0 ocr::fJoco(J(D 0= Z

·§.1.00 .....

7.51 U5 . CIl

0.75 ('oj

7.0"0

0

0.50 6.5

0.25 6.0 oooo~ =- ....... ............. --

0 5.5 0 2 4 2 4

10 log t (sec.)

Fig. 6.23. Creep and recovery of a PpPTA fiber (Twaron) for 0"0 = 1-1 GN m - 2:

creep strain (e) and dynamic compliance (0).208

The logarithmic dependence of the creep and the stress relaxation is characteristic of the response of a single-phase structure having a constant distribution of activation energies. Indeed, a single-phase structure has also been deduced from the SAXS, W AXS and ED studies of these fibers. It implies that the spectrum of the logarithmic retardation times L(ln ,) is independent of In, and the loss tangent is independent of the frequency.2IO,2I1 According to Ericksen this is caused by a broad range


1.04 6.06

1.00 .. 602 .. .......... ~ 0.96

...... --....... , 5.98 ·z ~ E

'" 6

5.94 :j> ~ 092 6 tJ./:;.M/:;.

~ Ui ~"'" • ""'6 0

0.88 lJ&..tJ.~l::. 5.90 ~6"""

0.84 5.86 0 2 3 4

10 log I(sec)

Fig. 6.24. Stress relaxation of a PpPTA fiber (Twaron) for £ = 1 %: stress ( ... ) and dynamic compliance (6).208

of intercrystalline bond distances and angles in the boundaries between crystallites.207

Assuming that the relaxation process with constant distribution of activation energies is confined to the shear deformation, a relaxation function for the orientation parameter u(t) = <sin2 cp(t) is introduced having the form

u(t)=uo exp[- ao -~ In (!..)] 90 90 to

(6.22)

where Uo is the value for <sin2 cp) before loading and J.l is a relaxation constant. 208 The logarithmic creep relation is now derived from eqns (6.18) and (6.22). For t = to it yields the relation for the change of <sin 2 cp) due to a true elastic extension of the fiber. Using eqn (6.16) it can be shown that for J.l/9o ~ 1 the creep rate of the primary creep is given by

(6.23)

According to this relation, highly oriented fibers should show a smaller creep rate than less-oriented fibers, which has indeed been observed. In addition, the rate of the primary creep should decrease for increasing stress. Actually it has been observed that for tensions larger than the yield stress the creep rate of mechanically conditioned fibers is constant and thus independent of the stress. The discrepancy with the prediction can again be understood by taking into account the effect of the finite width of the slipping element.


So far the results have shown that the viscoelastic behavior of paraaromatic polyamide fibers and of other well-oriented fibers can be described by a serial arrangement of an energy-elastic spring representing the chain modulus, and a viscoelastic element. The latter embodies the chain rotations due to shear deformation of the crystallites with a retardation time spectrum f(r) whose distribution in case of aramid fibers is described by the function r -1.208 This viscoelastic element may be regarded as a parallel configuration of an elastic spring go and a kind of dashpot having a stress relaxation function do/d(ln t) = - p. It acts like an elastic spring go at to, being the time at which creep or stress relaxation starts.

Dynamic mechanical measurements (DMA) provide information that is complementary to the creep and stress relaxation experiments. Frosini and Butta studied amorphous wholly aromatic polyamides and found for PpPT A small relaxation peaks at about 15De, 145°e and above 330°c.212 Except for the peak at 15°e, the peak heights are much smaller than those measured for aliphatic and partially aromatic polyamides. The peak at 15°e is identified as a fJ-relaxation and probably caused by the motions of free amide groups. The a-relaxation or glass transition peak was found at 320°C. Kunugi et al. performed dynamic mechanical measurements on PpPT A fibers spun from an isotropic solution and subsequently annealed at various temperatures up to 500°C. 213 They observed peaks with very low values for tan 15 «0'08): an a dispersion at 4600 e with an apparent activation energy of 770 kJ mol- 1 identified with the motion of the rigid backbone; a small fJ* peak at 270De arising from thermal molecular motion inside the crystallites; a fJ peak at 600 e with an activation energy of 206 kJ mol- 1 due to motions of the amide group; and a y peak near - 30°C. The fJ peak is very moisture sensitive and disappears almost completely in fully dried fibers. Approximately the same temperatures for the a and fJ relaxations were found by Badayev et al. 199

Figure 6.25 shows a dynamic mechanical curve of a PpPT A fiber measured in our laboratory. The relative flatness of the tan 15 curve is in agreement with the observation of the logarithmic creep. The low peak between 0 and 500 e is probably caused by a very small amount of water. As shown in Section 6.4.5, differential scanning calorimetry (DSC) indicates for a PpPT A polymer sample a glass transition temperature Tg near 290°C. Meaurements with Dse and DMA on PpPT A fiber samples do not reveal a Tg, which presumably is caused by the high degree of order and orientation. The DMA curve of the PBO fiber in Fig. 6.25

E z ~ UJ


200 -.----------------------TO.1

100

80

60

40 0.04

20

10+--.--,r--.--.--r-,--.--,r--.--.--.-+0 -200 -100 0 100 200 300 400

He)

329

Fig. 6.25. Dynamic mechanical analysis curves of PpPTA (Twaron HM) and PBO fibers, obtained in nitrogen atmosphere, frequency 72 Hz; filled symbols are

for PpPT A and open symbols for PBO.

shows that the decrease of the modulus with increasing temperature is smaller than for PpPT A fiber. This is consistent with the better thermal stability observed for PBO and PBT fibers (see Section 6.4.5).

Delayed elasticity is a property that is characteristic of a disorderly molecular arrangement in amorphous material or of disordered regions in crystalline material. 207 It is caused by thermally activated processes and may therefore involve entropy-elastic forces. Two observations suggest that the time-dependent elasticity of fibers made of para-aromatic polyamides, PBO and PBT, may have an entropic component. In the first place, the linear thermal expansion coefficient of these fibers is neg ative/95.214-216 measurements in our laboratory on a high-modulus PpPTA fiber in the range of -40 to llooe gave a constant value in this range of - 3·4 x 10 - 6 K - I. Secondly, the force on a fiber as a function of the temperature measured at constant length increases linearly with temperature, whereas in the case of energy elasticity this force decreases with temperature.214 As shown in Fig. 6.23, a broadening of the orientation distribution takes place during recovery from creep. This implies an increase in entropy, which presumably may be caused by the entropyelastic forces arising from the thermal fluctuations of the chain segments. So it is the viscoelastic 'shear spring' describing the chain rotations which provides the link between the creep strain and the time-dependent change in orientation.


6.4.3 Strength of Fibers and Films If the polymer chains in the fiber were as long as the gauge length of the fiber test specimen, and were oriented perfectly parallel to the fibre axis, the stress measured at fracture would be the theoretical failure stress. This stress, which is determined solely by the covalent forces in the chain, is estimated to be 10-15% of the theoretical chain modulus. 217 The calculation of the strength of real fibers and yarns is rather complicated owing to several factors such as the chain length distribution, the disorientation of the chains with respect to the fiber axis, the strong anisotropy of the forces between the atoms (covalent forces in the chain and secondary forces between the chains), the non-uniformity of the structure and morphology in the cross-section and along the fiber and, last but not least, the presence of impurities and voids.

The statistical description of the strength was introduced by Weibu11.218.219 In his model the assumption is made that failure is due to sudden catastrophic growth of pre-existing defects corresponding to local failure stresses. Failure at the most serious defect, i.e. the defect with the lowest fracture stress, leads to immediate failure of the fiber. It is further assumed that the defects are uniformly distributed throughout the fiber. The cumulative failure probability function, P, which represents the fraction of fibers that fail at or below a stress (J is, according to Wei bull, given by

(6.24)

where Lo is the test length of the fiber, CTo is a normalization factor and m is the Wei bull modulus describing the width of the strength distribution. In practice, the quantity In( -In(1- P)) -In Lo is plotted versus In CT, since this yields a linear relation of which the slope m is called the Weibull modulus:

In[ -In(1-P)] -In Lo =m In CT-m In CTo (6.25)

Figure 6.26 shows a Wei bull plot of a PpPT A yarn and illustrates the fact that the strength of fibers is governed by the effect of inhomogeneities and impurities.22o.221 A large Wei bull modulus implies a limited scatter in the fracture stress data, so weak fibers are absent. If the Weibull modulus for the fracture stress, as well as for the fracture strain, becomes very large (-100), the average fracture stress should approach the theoretical strength.

A theoretical study of the influence of the molecular weight and of the effect of the anisotropy between intrachain and interchain forces on the


6~--------------------~

-1 +------.-----.-----.-----1 3 3.5 4 4.5 5

Strength (GNm-2 )

Fig. 6.26. Weibull plot of the fracture stress distribution of a PpPTA (Twaron) yarn with m = 10·6.

fiber tenacity has been made by Termonia et al. 1S8 160 This study considers the case of perfect fibers made of an ordered array of fully extended and parallel-oriented chains, with no defects other than chain ends resulting from a finite molecular weight. It is based on the kinetic theory of fracture, in which the bond ruptures are simulated by a Monte Carlo process on a three-dimensional array of nodes. Strong bonds between the nodes in one dimension account for the covalent bonds in the chain and weak bonds between nodes in the other two dimensions represent the secondary forces. The results of this study are quite interesting and in many cases agree well with general observations on the strength of highly oriented fibers. For polyethylene fibers with low molecular weights, for example, intermolecular slippage involving rupture of secondary bonds occurs in preference to chain fracture, yielding tensile curves that at the end are bell-shaped. At high molecular weights primary as well as secondary bond rupture occurs, yielding tensile curves with brittle fracture. This gradual change in the failure process from slippage to chain fracture leads to the absence of a simple relation between fiber tenacity and chain length. In the case of PpPT A fibers the model shows that the process of fracture is initiated by the breaking of a small number of primary bonds and not by hydrogen-bond failure. This leads to a rapid build-up of stress concentration owing to the high concentration of chain ends resulting from the relatively short chains, and eventually to a brittle fracture of the fiber.


Apparently this model in its present form does not predict a fibrillated fracture morphology caused by shear failure, as has been clearly observed for the PpPT A, PBO and PBT fibers. Knoff noticed the similarity between the tensile failure morphology of PpPT A fibers and that of a uniaxially oriented fiber-reinforced composite. The latter fails in tension via matrix shear failure initiated at the fiber ends. This made him conclude that, if the shear forces at a discontinuity exceed the shear strength of the bond between the fibrils, the fiber tensile strength should be proportional to the fiber shear strength. 19B

Indeed, fracture morphology is probably the key for the understanding of the factors determining the fiber strength. Disregarding the effect of stress concentration due to impurities, we may expect that brittle fracture, characterized by a small fracture surface normal to the filament axis, is likely to be caused by the normal stress, (J n = (J cos2 cp, on the crystallite, and directed along the c axis, whereas a fibriIIated fracture surface is created by the shear stress, (J s = (J sin cp cos cpo The assumption is now made that primary bonds break at a stress equal to a specific fraction (10-20%) of the chain modulus, ee, and that the secondary bonds break at about the same fraction of the shear modulus, go. This implies for PpPT A a critical normal stress of about 30 GN m - 2 and a critical shear stress of about 0·3 GN m - 2. Evidently the latter value is easily attained during tensile testing of a medium-oriented fiber. Since for increasing orientation it becomes less probable that there will be a considerable fraction of chains oriented at an angle for which the critical shear stress is reached, this approach leads to an increase of the strength with increasing modulus. In highly oriented fibers, having a rather contracted orientation distribution, the overall shear stress will not reach the critical shear stress. Therefore, the observed shear failure in these fibers is presumably caused by impurities or inhomogeneities causing shear stress concentrations, since in structures with a strong anisotropy of the modulus (e.g. primary versus secondary forces) large shear stress concentrations arise near a crack tip. These shear stresses are parallel to the direction of the largest modulus value. 217 Thus, for highly oriented fibers containing impurities or inhomogeneities the strength will hardly depend on the fiber modulus.

An interesting and rather simple model for the fiber strength has been developed by Yoon. 222 It uses the same principles that have been applied to estimate the strength of short-fiber-reinforced composites, i.e. an elastic load transfer with a debonding matrix which occurs through elastic deformation of the interface. In Y oon's model the fiber is not regarded as a continuum, but as a composite consisting of covalently bonded chains


in a matrix, whose properties are solely determined by the weak secondary bonds. It is based on the consideration that the macroscopic load is transferred by interchain interactions and that rupture of the fiber occurs when the shear force on the chains exceeds a critical value. Then the chains are separated and fiber failure is caused by the rapid accumulation of the interchain voids, resulting in the observed fibrillar fracture. The model also takes into account the orientation distribution of the chains. Good agreement with experimental results was obtained for wholly aromatic polyesters.

An alternative approach to the fracture process deals with the effects of cracks. It considers the balance between th" ~tored elastic energy per unit volume that will be released upon crack propagation and the work of fracture per unit area of crack surface. Cracks with dimensions smaller than the critical Griffith length, L G , are stable, whereas cracks larger than LG are self_propagating. 223 ,224 As has been shown previously, paraaromatic polyamide fibers, PBT and PBO fibers, have many elongated voids parallel to the fiber axis. An analysis of their effect on the strength following the Griffith approach has not yet been given.

A detailed study of the strength and lifetime under constant stress of single PpPTA filaments using Weibull statistics and an exponential kinetic breakdown model was carried out by Wu et al. 225 They found that filaments failed owing to transverse crack propagation after very short creep times, but that after long creep times the failure mechanism was splitting and fibrillation. Activation energies of the failure process amounted to 340 kJ mol- 1, which seems to indicate rupture of the C~N bond in the chain backbone.

As mentioned in Section 6.4.1, a heat treatment under tension of as-spun PBO and PBT fibers increases the modulus as well as the strength substantially.201,202 As with the aromatic polyamide fibers, the enhancement of the fiber modulus can be explained solely by the contraction of the chain orientation distribution. But no clear explanation is given for the improvement of the strength, although it has been suggested that this is due to the improved lateral order. An alternative explanation is the following. The heat treatment at elevated temperatures increases the conjugation in the PBT chain. 115 This causes a decrease of the rotation angle between the phenylene and the benzobisthiazole segments, which in turn yields a smaller cross-section of the chain and thus a larger chain modulus. As shown by eqn (6.15) this phenomenon, together with the contraction of the chain orientation distribution, can have a large effect on the fiber modulus. Assuming that fracture is only


caused by the shear component of the applied stress, the strength of fibers containing hardly any impurities and structural imperfections is expected to increase with the fiber modulus, as has been discussed before.

The strength of films is strongly influenced by the kind of texture, as has been demonstrated by Haraguchi et al. 117 They made PpPT A films with a (OkO) uniplanar orientation (hydrogen-bonded planes oriented perpendicular to the film surface) and with a (hOO) uniplanar orientation (hydrogen-bonded planes lying parallel to the film surface) using water or an organic liquid as coagulant. The (OkO)-oriented films showed a ductile tensile curve with a strength two times higher and an elongation at break five times larger than the values obtained for the (hOO)-oriented films. This difference is explained by assuming that the hydrogen-bonded planes play a primary role as slip planes during tensile deformation of films.

Film preparation by different techniques, such as extrusion, uniaxial draw-down and mandrel processing, has been explored by Aoki et al. 226

and Bodaghi et al. 227 The uniaxially oriented PpPTA films obtained by the drawing process exhibit highly anisotropic mechanical properties, whereas the mandrel-produced films show balanced properties. However, they found that the values for modulus and strength of these films are much lower than those obtained for fibers.

This is also true for PBT, as is shown by the study of as-extruded and heat-treated films of Feldman et al. 228 The maximum values obtained for modulus and strength, 238 GN m -2 and 1·51 GN m -2, respectively, are considerably lower than the largest fiber values (see Table 6.3). A study of the temperature and strain rate dependence of the deformation behavior of these films revealed the onset of a structural reorganization near 300°C, while the stress activation volume characterizing the activated rate process of the yield stress increased considerably above 200°C. 186

6.4.4 Compressive Strength of Fibers F or application as reinforcing components in composites the fibers should have not only a high modulus but also a certain level of compressive strength. Greenwood and Rose were the first to investigate the compressive behavior of high-modulus PpPT A fibers in unidirectional composites by means of the elastic loop test. 229 At about 0·8 GN m - 2 compressive yielding accompanied by the formation of kink bands was observed in the filament. Kink bands, revealed as dark lines in the polarizing microscope, have also been observed in PpPT A fibers which had been subjected to a tensile fatigue test. 230 The kink bands are associated with well-defined regions in which there is an abrupt change of


the chain orientation with respect to the fiber axis.231 These shear bands may extend across the filament diameter. Application of a tensile stress shows that the induced kink bands are capable of realigning along the fiber axis.

The kink and slip bands formed at compressive yielding are largescale manifestations resulting from the buckling of chains. At locations where buckling takes place the chains are forced to adopt a very different conformation which may extend over a large part of the chain. Largescale segmental movements are also associated with the glass transition temperature of a polymer. The apparently common origin of phenomena such as compressive yielding and the glass transition, related to chain flexibility and intermolecular interactions, suggests a correlation between the compressive strength, (le, and Tg• Indeed an empirical relation has been found, viz. (le is proportional to T: for materials composed of mainly first- and second-row atoms. 232 This result can also be derived quantitatively by assuming that the work for compressive yielding is proportional to the activation energy of the glass transition.

Deformation band studies of axially compressed PpPT A filaments revealed two distinct types of kink bands. In filaments that show tangential splitting, bands at an angle of about 55° are observed, whereas in radially split fibers the kink bands are oriented perpendicular to the fiber axis. Hydrogen-bonded planes acting as slip planes and intermicrofibril slip play an important role in the deformation process during axial compression.233

Tensile tests of axially compressed fibers revealed only a 10% loss in tensile strength, after application of as much as 3 % compressive strain. 197

In this study the regular spacing of the helical kink bands at 50° to 60° to the fiber axis was noticed.

As shown in Table 6.4, the compressive strength of aromatic polyamides, PBO and PBT fibers is considerably higher than that of aromatic polyester or ultra-high-molecular-weight (UHMW) polyethylene fibers.234-236 The compressive strength of PpPTA fibers increases with increasing modulus, as is shown in Fig. 6.27, and the critical strain for kink band formation decreases continuously from about 1 % for a low-modulus fiber to about 0·5% for a high-modulus fiber. It was also observed by Van der Zwaag that the kink bands are formed before elastic instability occurs, so that they may therefore be attributed to a plastic deformation process. Furthermore, a model has been developed that describes the density of kink bands as a function of the compressive strain235 ,236 (see Fig. 6.28).


Table 6.4 Compressive strength of some polymer fibers

Fiber

UHWM PE Aromatic polyester PpPT A low modulus PpPT A high modulus PBO PBT

Compressive strength (M N m - 2)

50 200 500 820 680 680

~E 10 -y--------------, z S2 :5 0.75 rn c ~ (j) 0.5 c Q

"' "' ~ 0.25 a. E o U O+---.---.--.--~

o 50 100 150 200 E-modulus (GNm-')

Fig. 6.27. Compressive strength of PpPT A fibers versus the elastic modulus.235

150.------------.

100

50

0-+-......... --&..---. ...

o -0.5 -1.0 -1.5 E(%)

Fig. 6.28. Kink band density in a PpPT A filament versus compressive strain. The model yields the solid curve.236


The results for the aramid, aromatic polyester and PE fibers reveal that the values for the compressive strength can differ widely between different polymeric fibers, notwithstanding comparable tensile properties. This is due to the combined effects of differences in bending stiffness of the chains and differences in lateral cohesive binding forces. In the case of the PpPT A fibers intrinsically stiff chains are bounded by a regular network of hydrogen bonds, whereas in polyethylene fibers the flexible chains are only weakly bonded by Van der Waals forces. With regard to the effect of the interchain forces it should be noted that, though the chains of PBO and PBT are less flexible and the fibers have a higher modulus than the high-modulus PpPT A fibers, their compressive strength values are lower.234.237 Presumably the hydrogen bonds in PpPTA account for this difference in compressive properties. Glass transition temperature as well as compressive strength are both determined by the combined effect of chain properties and intermolecular forces.

Compared with glass and carbon fibers the compressive strength of fibers made from lyotropic polymers is rather low. However, as shown by Van Dreumel, the negative effect of a low compressive strength on composite properties should not be overrated.238.239

6.4.5 Thermal Properties Owing to the rigid molecular chains and the high crystallinity, PpPT A and the other para-aromatic polyamide fibers show high thermal stability. The differential thermal analysis (DT A) curves as well as the differential scanning calorimetry curves (DSC) of the PpPT A polymer and fiber show an endothermic peak between 500 and 600°C (see Fig. 6.29). This peak is only observed at high heating rates and is superimposed on the slope of a large decomposition endotherm (with a maximum temperature above 600°C). The peak temperature is independent of the heating rate, thus confirming the observations of Brown and Ennis and their interpretation of this phenomenon as a melting transition.240 The onset of decomposition, on the contrary, is strongly dependent on heating rate. Depending on the fiber or kind of polymer, the peak temperature varies between 550 and 565°C. The reversible nature of the melting process has been confirmed for the polymer samples. 241 The onset of the degradation endotherm corresponds with the onset of weight loss in the thermogravimetry (TG) curves. At a heating rate of 20 K min -1 the TG curve of a PpPT A (Twaron) fiber shows a weight loss of 2·5% at 550°C. These results are in general agreement with earlier investigations.242-246


40,---------------------------------------~~

5; E

30

-; 20 o u: rn __ ---1. 50. ,--,--

10 _,--

~

J.- -- -- ---,/' /,/ '560~c:; .. ,,"·"···

/'

/ /'

/'

--" 20.0. __ ---'--' O+-~._--_r--~--~--~----~--~--~--~--~

500 550 T('C)

600

Fig. 6.29. DSC curves of PpPT A polymer at various heating rates obtained in nitrogen atmosphere, showing a melting phenomenon. 241

As shown in Fig. 6.30, a DSC curve of dried PpPT A polymer shows a small endothermic baseline shift, in conjunction with an endothermic peak near 290°C, probably representing the glass transition temperature of the polymer. This Tg value agrees with the one found by Brown and Ennis.24a It is also in line with the observation found for a wide variety of polymers that Tg/Tm =0·66. However, it differs from the values found

2.4,------------------------------------------"

2.0

~. 1.6 E

;: 1.2 o u:

'" -1. 0.8

0.4

Tg

287'C

O+---._--_r---r---.---,--~._--r_--~--~--~

250 300 T(CC)

350

Fig. 6.30. DSC curve of PpPT A polymer at a heating rate of 100 K min - I.

obtained in nitrogen atmosphere, showing glass transition.241


with dynamic mechanical analysis (see Section 6.4.2). Annealing of PpPT A fibers showed that at temperatures up to 350°C the crystallinity remained constant. Near 400°C it reached a maximum, and decreased at temperatures exceeding 500°C. Prolonged thermal aging at 150°C showed some increase in lateral crystallite sizes.247

Prolonged exposure of PpPT A fibers at temperatures up to 250°C for more than 100 hours results in a relatively small decrease (about 17%) of the strength. Short-time exposures up to 450°C can be applied without large strength losses. The temperature dependence of the strength in the range from about -40°C and higher is much smaller than that of conventional synthetic fibers such as PETP and the nylons. 248

PBT fibers show a very small strength loss of only 2% after an exposure of 65 hours in air at 300°C. This indicates a thermal stability of the mechanical properties which is even better than that of the aromatic polyamide fibers.31.183

Heat treatment at 600°C of as-spun PBO fibers increases the modulus by almost a factor of 2, but the strength is also slightly improved (see Table 6.4). Higher temperatures cause a reduction of strength due to degradation. 143

Figure 6.31 demonstrates the effect of the temperature on the Young's modulus of PpPT A fibers with different initial moduli. 249 For decreasing temperatures a considerable increase of the modulus is observed, which presumably is mainly due to an increase of the modulus for shear, go, between adjacent chains. As ti1etemperature drops, this increase is caused by the stiffening of the intermolecular bonds. To a lesser extent the decrease of the orientation parameter, <sin2 q», may have some effect (see eqn (6.15)). Like the negative thermal expansion coefficient observed for PpPT A fibers (see Section 6.4.2), this expected decrease of the orientation parameter is of entropic origin.

Creep measurements of PRD 49 (PpBA) fibers at 65°C and 150°C exhibit logarithmic creep, as is also observed for PpPT A and PpBAT fibers at room temperature, and the creep rate increases with temperature. 207 Wu et al. determined Weibull parameters of the strength and lifetime distributions of PpPT A fibers at room and elevated temperatures. 225 They found that the mean filament strength (and Weibull scale parameter) varies inversely with temperature, while the strength variability (and Weibull modulus) remains practically constant over the range of temperatures used (20-130°C). Creep-rupture lifetimes showed very high variability (small values of the Weibull modulus), but at 130°C the Wei bull modulus increased with decreasing stress. The activation energy derived from the lifetime measurements amounts to 330 kJ mol- 1 and


160 -., . ..........

"'-. """. 140 -. '" ...... .", ..........

"... .~ ~ •

~ 120 --........... " "'. .~ "'-~ .......... .~ .,

<Jl "'-. ~ :::l """. •

"0 ~. 0

'" ~ 100 ~ w ."'" , .",

• 80 "'-.

" • ,

-200 -150 -100 -50 0 He)

50 100

Fig. 6.31. Modulus versus temperature for PpPT A fibers with different Young's moduli. 249

indicates that chain scission is the dominant mechanism in creep-rupture experiments.

ACKNOWLEDGMENTS

We thank F.A.M. Schenkels, H.G. Weijland, 1.1. van Aartsen, S. Picken, S. van der Zwaag, H. Jansen, A. Weed a and H.M. Heuvel for critical reading of the manuscript and many helpful discussions. In particular, the technical assistance of Mrs B. Schaffers-Korff is gratefully acknowledged.

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Chapter 7

Thermotropic Side Chain Liquid Crystal Polymers

Derek J. Simmonds Division of Chemistry, School of Science, Sheffield Hallam University,

Pond Street, Sheffield, UK, S1 1 W B

7.1 INTRODUCTION

A simplified representation of a polymer containing lath-like, side chain substituents is shown in Fig. 7.1(a) (cf. Table 1.1 in Chapter 1); the term 'comb polymer' is an attractive description for this type of structure. Each repeating unit contributes a short length of the comb's spine and one tooth. As long as the motions of the polymer backbone do not unduly restrict the freedom of meso genic teeth to form ordered assemblies, then we might expect such materials to reproduce, to some degree at least, any liquid crystal (LC) behaviour associated with the mesogen units as small-molecule liquid crystals (SMLCs). We have then a model for comb liquid crystal polymers (comb LCP) or side chain liquid crystal polymers (SCLCPS).l

TIl 0 0 (a)

Fig. 7.1. Side chain liquid crystal polymer (a) without and (b) with a flexible space.

349

350 Derek J. Simmonds

In fact our simple model is both unduly restnctlve and misleading. Unless the mesogens are separated from the backbone via flexible 'spacer groups', e.g. (CH 2)n, then they can have little independent mobility. Since flexible polymer backbones tend to adopt random coil conformations, long-range ordering of mesogens is unlikely in most cases unless appropriate spacers are incorporated into the structure to decouple, at least partially, the motions of meso gens and backbone. Attention to the spacer conceptl has led to the synthesis of a wealth of comb polymers that exhibit LC phases. An improved representation is therefore Fig. 7.1 (b), while Warner3 and others4 have proposed more realistic models whereby a random coil backbone can support mesogens that, owing to the presence of flexible spacers, are able to adopt ordered dispositions relative to some director vector ii (e.g. the result of an electric or magnetic field), as shown in Fig. 7.2. Such LC materials could have interesting and applicable properties as detailed at the end of this chapter.

(a)

(b)

Fig. 7.2. (a) Alignment parallel with director. (b) Alignment perpendicular to director.

Thermotropic Side Chain Liquid Crystal Polymers 351

7.1.1 Historical Development and Literature With hindsight we can discern stages in comb LCP development. Demonstration of the successful exploitation of the spacer concept represents a key link between earlier syntheses and the intense activity that was to follow. By the mid-1980s a vast range of polymers had been produced and a further development was detectable; strategic molecular design and structure-property considerations were increasingly superimposed upon synthetic efforts. Meanwhile, previously unexploited backbones and spacers, and new meso gens, were emerging in the literature. Ample commentary on the maturing science of LCPs was provided by various books5, while detailed reviews6 provided comprehensive summaries of what had been achieved and set out to identify important parameters leading towards a rationalization of structure-property relationships.

In parallel with the synthetic developments, work moved ahead in applying analytical methods to the modelling and phase investigation of the new comb polymers, and possible areas for application were investigated; comb LCPs became a field of study in their own right. 7 Progress in materials science and technology requires the constructive interaction of scientists from several disciplines to interweave, for example, synthesis, characterization, evaluation and application into a progressive sequence propagating via results, review and feedback. All the various threads of activity up to 1988 were drawn together by McArdle! into an encyclopaedic collection of contributions from influential workers involved with side-chain LCPs (SCLCPs). The result summarizes earlier work and sets out an agenda for the next stage, which is increasingly based on the application and exploitation of SCLCP products.

7.1.2 Scope and Nomenclature We are chiefly concerned with LC structures wherein mesogenic entities are present as substituents connected (generally via a spacer) to a flexible polymer backbone that is not itself inherently meso genic, although combined or double systems (Table 1.1 in Chapter 1) will be mentioned. The mesogens confer the property of mesophase formation (i.e. between conventional solid and liquid phases) upon the composite structures under thermal control (hence 'thermotropic').

Referring to the nomenclature defined in Table 1.1, we shall encounter mainly one-comb, double-comb, parallel (and biparallel) and network LCPs that exhibit calamitic phases (such as smectic and nematic). Calami tic phases result from mesogens that are lath-shaped. We shall


survey structural and synthetic aspects of these systems and note some potential applications. As we shall see, copolymers based on the listed designs but containing more than one type of substituent are becoming increasingly relevant to the future of comb-like systems; the palisadecomb (Table 1.1) is one possible outcome of the copolymerization approach. Disc-combs involve mesogens that are disc-shaped rather than lath-shaped, and disco tic phases may result from appropriate polymer designs; disc-combs, like double polymers, receive only brief coverage here.

The important LC phases have been described in Chapters 1 and 2. Mesophase characterization typically requires the combination of thermal analysis-e.g. differential scanning calorimetry (DSC)-to establish transition temperatures, with hot-stage microscopy and/or X-ray diffraction to identify the mesophases. Since we are dealing with polymeric systems we must remember the importance of thermal history (e.g. whether annealed or not annealed), and heating and cooling traces (DSC) are frequently non-identical. Monotropic mesophases are quite common; these are metastable phases that are detectable only in a cooling experiment from an isotropic state. They are generally associated with polymers that display a melting transition and are obtained on supercooling. By contrast, enantiotropic phases are thermodynamically stable and can be observed between a melting point, or glass-transition temperature, and a clearing point whether the sample is being heated or cooled.

Thermal data quoted in this chapter present transition temperatures between phases that are abbreviated as follows:

g = glassy phase, lost on heating beyond the glass transition temperature, Tg

K = crystalline phase S = smectic mesophase (e.g. SA, Sc, etc., and S~ for a chiral smectic

C phase) N = nematic mesophase (RN being a re-entrant nematic phase)

Ch = cholesteric mesophase (chiral nematic) i = isotropic liquid phase, fully realized above the clearing tem

perature Tel.

In this chapter, structure--property correlations relate to thermal transitions almost exclusively although enthalpy and entropy changes are also reported where appropriate. However, other parameters are of considerable importance; for example order parameters (indicating the deviation from complete alignment with a director) are clearly relevant


even though we shall not encounter them here. Birefringence, which describes the visual effectiveness of an LC in display applications, results from different refractive indices of LC fluids parallel and perpendicular to their alignment direction. It is not considered further in this chapter; suffice it to say that the mesogens we shall encounter all contain conjugated double bonds and conjugation is an important contributor to satisfactory birefringence.

Phase transition temperatures can help define operating limits for LC materials, but the availability of a desired meso phase at a chosen temperature is still only part of the story as far as potential applications are concerned. The speed and efficiency of alignment or realignment are also relevant. The intrinsic viscosity of a material will have an effect here, and polymer LCs are much more viscous than SMLCs, which rules out applications that require very rapid alignment switching as far as polymers are concerned. Mesogens themselves are associated with dielectric constants and elastic constants that help define their response to aligning influences and their subsequent relaxation to an equilibrium ordering.s

As covered in Chapter 4, dielectric constants (c) may be measured parallel or perpendicular to an applied field (c and C.1, respectively), and the difference between those values (~c=c; -C.1) is of major importance. A positive ~c implies alignment parallel to an applied field (homeotropic) while perpendicular (normal) alignment results if ~c is negative. The magnitude of dielectric constapts reflects the dipole moment of the mesogenic entity and therefore the nature and position of substituents on the mesogen core are of significance with regard to s as well as to thermal parameters (see below). Dielectric relaxation spectroscopy9 is a useful technique for probing dielectric alignment phenomena in comb LCP evaluation, and is discussed in Chapter 4.

Elastic constants (k) reflect the flexibility of mesogens to change their alignment via processes described as either splay-mode (k 11 ), twist-mode (kzz ) or bend-mode (k33) distortions from an existing situation. A key parameter (k33/kl1) needs to have a low value for efficient display properties. Mesogens containing aromatic rings are generally satisfactory in calami tic systems and all the mesogens considered here are aryl species.

7.2 GENERAL STRUCTURAL FEATURES

An effective comb LCP contains three primary structural variables each of which is significant to a realization of desired thermotropic behaviour,


processability and end-use. Each is also further divisible into additional variable components; hence the complexity of Section 7.3 where we seek to identify correlations between structural and phenomenological variations. The three primary variables are the polymer backbone, the spacer linkage and the mesogen itself. Having noted the more important constituents relevant to comb LCPs with lath-shaped mesogens, we look briefly at other classes of side chain system in Section 7.2.4.

7.2.1 Polymer Backbone The backbones encountered are generally familiar in thermoplastics technology and tend to be flexible polymers associated with glassy, rubbery or fluid phases depending on temperature. Overwhelmingly dominant are polymethacrylates, polyacrylates and polymethylsiloxanes, in order of increasing chain flexibility (decreasing Tg). These and other comb LCP backbones are listed, with indicative references, in Table 7.1.

Table 7.1 Backbones for side chain LCPs

Acrylics (acrylate, methacrylate, chloroacrylate)IO Siloxanes II and cyc1osiloxanes 12 Polystyrenes, polyvinylethers, polyoxiranes 10 Mesogen-containing backbones I 3 (for double LCP)IO Polyalkenes 14 Polynitriles 15 Polyacrylamides 16 Polyphosphazenesl 7

Polyurethanes I B

Polymalonates l9 Poly(vinylether-alt-maleic anhydride)20

Below mesophase temperatures backbone crystallization is not particularly desirable in LCPs (nor is it very common), but the stiffness associated with operating temperatures below 1'g can be used as a means of quenching in mesophase order attained at higher temperatures. Usually this means heating to isotropic fluid, slowly cooling to mesophase in the presence of an aligning field for macroscopic ordering, then rapidly quenching to below Tg to lock in the ordered arrangement (vitrification).

As well as the constitution of the backbone, it is necessary to acknowledge the influence of details such as tacticity, polydispersity and molecular mass (e.g. average degree of polymerization, DP). Finally we will look


at some polymers made with more than one type of substituent on the backbone, embracing copolymers and cross-linked networks; these will not, however, have more than one type of backbone in their structure.

7.2.2 Spacer linkage Although a few comb LCPs having direct attachment of the mesogens onto the backbone are known (Pugh and Percec have listed most of them21 ), the majority contain a flexible spacer linkage. The spacer provides partial decoupling of the motions of the backbone and mesogens, and contributes structurally to both the polymer chain (e.g. as an internal plasticizer)lO and the mesogen (e.g. as a tail group, cf. Section 7.3.3.1). It is therefore somewhat artificial to separate out the spacer, but structure-property relationships are easier to identify by doing so. Furthermore we can subdivide the total structure between the lath of the mesogen and the backbone of the polymer into the spacer chain itself and two connections, which are often functional groups. Some representative components are illustrated in Fig. 7.3.

Backbone Connection

Flexible spacer

Mesogen connection

Ester, Direct (Alkyl), Ether, Amide 16

I o I

CH2

I CH2 n

I Ester, Alkyl, Ether, Amide16, Carbonate22

Fig. 7.3. Examples of linkage components.6 •10

The most popular flexible spacer is oligomethylene (n = 3-11) but oxyethylene and siloxane spacers are not uncommon. The identities of the connecting groups, like that of the spacer itself, have implications for the flexibility of the total side chain. More subtly, the connection onto the mesogen may extend its rigid lath structure; the connection -O-CO-aryl extends the n-system of an aryl mesogen while -CO-O-aryl does not.

A further variant on the spacer-mesogen linkage follows from the lath shape of typical mesogens. In most comb LCPs the linkage is to the end of the long axis of the mesogen (longitudinal attachment), but connection


to a point about half-way along the lath (lateral attachment) produces parallel (or biparallel) LCPs. Examples of both modes of attachment appear in Section 7.3.2.

7.2.3 Mesogen We are concentrating on lath-shaped mesogens (see next section for examples of discotic mesogens). These contain three significant components: cyclic structures (frequently benzene rings), groups linking the rings together, and substituents on the rings. On the whole birefringence is related to the degree of conjugation in the mesogen, while dielectric anisotropy (~£) depends on the nature and position of substituents. Figure 7.4 is a sort of menu of typical components in a mesogen containing two benzene rings with substituents on the longitudinal sites only (the most common type of structure).

Spacer ---O-x---O-y o 0

e.g. II II X=direct linkage, -C-O-, -O-C-, -N~N-, -CH=N-

e.g. Y =CN, alkyl, O-alkyl, NO z

Fig. 7.4. Common mesogen components. to

Comb LCPs with efficient decoupling via the spacer are, in principle at least, polymer-bound mesogens. We may then expect close parallels between the thermal profiles of LCP and the SMLC of the same mesogen. This is only partly borne out in practice, and a 'polymer effect' is now widely acknowledged; incorporation of a mesogen into a LCP leads to meso phases of higher order. Thus, for example, a mesogen that is nematogenic as a SMLC may produce a smectic polymer, and indeed appropriate structures that are not actually liquid-crystalline as monomer species may produce nematic polymers. Several examples of comparative data for SMLC/LCP pairs are given in Ref. 23.

Chiral mesogens generate chiral meso phases, in which successive layers of meso gens are aligned such that the dipole moment vectors for each layer are somewhat displaced from those of the neighbouring layers, giving a helical twist of the vectors if viewed through sufficient layers. When the ordering in each layer is nematic, then the resulting chiral nematic phases are often called cholesteric (after cholesterol, a steroid


nematogen). Cholesteric phases exhibit selective reflection of light of different wavelengths, a property of use for thermochromic materials. 24

Certain cholesteric LCs also give rise to the so-called blue phases.24

Chiral smectogens may generate chiral smectic phases of which the Sc (chiral smectic C) phase is of great interest. Sc is a tilted smectic phase and Sc is additionally subject to helical periodicity of the director in successive tilted layers. However, the symmetry properties of the tilted smectic phase imply a non-vanishing polarization in an Sc material when the helix is unwound, producing a ferroelectric material. Ferroelectrics have potential electro-optical applications and may be superior to nematic systems. 2 5

7.2.4 Other Variables: Discotic and Double Systems In a 'disc-comb' (Chapter 1, Table 1.1) the lath-shaped mesogens we have been considering are replaced by disc-shaped ones; if lath meso gens are pseudo-rectangular then discs are pseudo-circular. Figure 7.5 shows two such structures. Siloxane combs containing triphenylene (I) mesogens are liquid crystalline26 whereas acrylate combs with phthalocyanine (II) substituents have not yet given mesophases despite success with SMLC phthalocyanines,z7 Discotic ordering is illustrated in Fig. 1.7 (Chapter 1).

Computer modelling results28 raise the intriguing possibility that cyclosiloxane backbones (Table 7.1) substituted with calami tic mesogens could behave as discotic LCs or nematic LCs depending on the siloxane ring size and the spacer chain length; small ring sizes could allow an approximately planar arrangement of the mesogen substituents (III in Fig. 7.5), with disco tic stacking of the molecules a possible outcome. They would be small-molecule discotics rather than oligomeric side chain systems.

The 'double' architectures shown in Table 1.1 correspond to LCPs more commonly referred to as 'combined', indicating that their structures combine the features of main chain and side chain LCPS. 29 Thus meso genic entities are located both in the side chains and in the polymer backbone, and mesophase microstructures tend to contain mesogens from each of these environments. Ordering that involves main chain components implies a tendency away from the randomcoil conformation associated with flexible, non-mesogenic backbones; this is possibly why double LCP materials are often crystalline and tend to exhibit at least one smectic mesophase (sometimes several) in their thermal profiles.

358

Spacer I o

OR

Derek J. Simmonds

c~H-spacer

R0:OON~NWOR '7 ::;.-- NH HN I ~ ~ ::-... OR

RO N~N

>=< RO OR

III

Fig. 7.5. Possible discotic mesogens.

Most double LCPs are polymalonates prepared by the polycondensation of an 'alkyl' malonic acid (containing the side chain mesogen in the alkyl group) with a mesogenic diol, to incorporate mesogens into the main chain. Some examples are given in Section 7.5.3. Pugh and Percec have recently reviewed combined systems,10 and they subdivide the class into various groups. The substituent spacermesogen moieties can be attached to the main chain at points on its mesogens or from the spacers between the mesogens (e.g. the malonate systems). If the former, then the main chains can either include or not include flexible spacer units. Finally, cruciform mesogens incorporated into the main chain of a polymer, generating the class 'star' LCPs (or 'cross' LCPs) shown in Table 1.1, are also


combined materials, albeit without spacer groups between the main chain and side chain components; several of these materials were reported by Berg and Ringsdorf.30

7.3 STRUCTURE-PROPERTY CORRELATIONS

With such a varied menu of comb LCP components to choose from, it is important for synthetic chemists to appreciate how decisions about molecular architecture might influence the properties of polymer products. While the spacer, the mesogen and the constitution of the backbone are fixed by preselection of reactants, the average molecular mass, the polydispersity, the degree of substitution along the backbone, and the purity are variable. Systematic syntheses with these latter variables under careful control have been performed only quite recently. We shall review some of the results once we have covered the actual constitution of the backbone. A good deal is known about the effects of variations to the spacer and mesogen components, and the major correlations are listed in Sections 7.3.2 and 7.3.3.

On the whole, the correlations that are made concern structure versus phase transition phenomena. In particular we shall concentrate on the identity of mesophases that are formed, the persistence of the mesophase region (thermal stability of the mesophase) and the glass-transition and clearing-point temperatures (Tg and Ted. But other properties are important also and we shall look at a few correlations between structure and, for example, thermodynamic quantities and dielectric relaxation phenomena.

7.3.1 The Backbone We have already mentioned the polymer effect; a small-molecule mesogen is liable to produce more ordered phases when attached to any polymer backbone as long as sufficient attention is paid to molecular design (e.g. spacer concept). Here we are concerned with the effects of attaching the same mesogen-spacer combination to different backbones, the variables being the atomic constitution of the polymer, its molecular mass, polydispersity and tacticity. Unfortunately, no single mesogen-spacer combination has been subjected to all the variables in well-controlled experiments, so we must look at several systematic studies that involve different meso gens and spacers.


7.3.1.1 Constitution of the backbone In the context of comb LCPs the most important property of a polymer backbone is its flexibility. Finkelmann has pointed out31 that while different backbones may not cause different mesophases, they do at least bring about different transition temperatures. The major correlations are between backbone flexibility on the one hand, and Tg and mesophase persistence on the other. Tg decreases as backbone flexibility increases, while AT ( = Tel - Tg ) increases with increasing flexibility, although clearing temperatures themselves (the onset of I, the isotropic phase) do not necessarily increase with flexibility;-that depends largely on the length of the spacer, as we shall see later.

In Table 7.2 phase transition data are presented for three series of mesogen-spacer combinations attached to different backbones, showing the generality of the Tg and A T correlations for nematic methoxybenzoate combs, smectic cyanobiphenyl and nematic parallel side chain systems.

In Table 7.2 progress from left to right corresponds to increasing backbone flexibility, although chloroacrylates and methacrylates (first two columns) are of similar flexibility while cyclosiloxanes and linear siloxanes differ in terms of degrees of freedom rather than flexibility per se. In the cyanobiphenyl series the 'LC' phase (cyclosiloxane example) is probably smectic A but may be discotic 12b; note also an additional, nematic phase in the acrylate entry. In the final, parallel, series the siloxane entry actually has a slightly different mode of attachment of spacer to mesogen (ester) than that shown, so that the absolute data quoted are subject to error in direct comparisons with the other three entries. Even so the Tg and A T correlations are well seen, and they seem to be general in other, similar studies also.

7.3.1.2 Molecular mass and polydispersity Molecular mass averages, e.g. number-average (Mn) and weight-average (Mw ), are significant variables in polymerization processes. Furthermore the distribution of molecular masses, as defined by the polydispersity (Mw/Mn), may differ even if two polymers have one molecular mass average the same. It is not particularly easy or convenient to work with well-defined mass averages and polydispersities but, as the variations in this section show, it is quite important to do so. We shall look at some recent studies that throw light on relationships with varying polydispersities and varying molecular mass averages. In Tables 7.3(a} and 7.5 the molecular mass parameter is the average degree of polymerization (DP).

Tab

le 7

.2

Infl

uenc

e o

f ba

ckbo

ne o

n p

hase

tra

nsit

ion

tem

per

atu

re (

K)

and

mes

opha

se s

tabi

lity

(~T)

Cl

Me

H

CD

M

e I

I I

o-t

I B

ackb

one:

+

CH

2-

C+

+

CH2

-C-+

+

CH

2-C

+

+O

-S

i+

Spa

cer/

mes

ogen

~

I I

I I

00

-C

OO

-C

OO

-Q

(CH

2)2

0@

CO

O@

-O

Me

g 36

9 N

394

i g

32

0N

350

i g

288

N 3

34

~T=25

~T=30

~T=46

[Ref

. 31

J [R

ef.

31J

[Ref

. 31

J

(CH

2)5

0@

-@

-C

N

g 33

3 SA

393

i

g 30

8 SA

393

N 3

97 i

g 29

0 L

C 4

23 i

g

287

SA 4

43 i

~T=60

~T=89

~T= 1

33

~T= 1

56

[Ref

. IO

J [R

ef.

IOJ

[Ref

. 12

bJ

[Ref

. IO

J

(CH

2)"

C4

H9O

@C

OO

@-O

OC

-@

-O

C4

H9

g 29

3 N

339

i

g 29

2 N

34

0i

g 27

7 N

340

i g

290

N 3

61 i

a ~T=46

~T=48

~T=63

~T=71

[Ref

. 32

J [R

ef.

33J

[Ref

. 32

J [R

ef.3

4J

aDif

fere

nt s

pacer-

see t

ext.


Four general points emerge from studies by Shibaev,35 on acrylic comb LCPs. Phase transition data were compared for an unfractionated specimen and a series of carefully fractionated samples of different DPs. The four major conclusions were as follows:

- Most phase transition temperatures tend to rise with increasing DP. - There is a critical DP beyond which the dependence of Tel on DP

levels off. - The mesophase types observed may depend upon DP (e.g. a SF-N

transition was only observed at lower DPs). - Some mesophases exhibited by fractionated polymers (narrow dis

tribution of DP) may not be observed in unfractionated specimens even when the same DP is present within the broad distribution of the unfractionated sample.

Other studies support and extend these conclusions.36 The data shown in Table 7.3 (a and b)ll derive from carefully fractionated siloxane materials and clearly illustrate dependencies of both Tg and Tel upon DP (in Table 7.3(a), note the almost constant polydispersity), and also upon polydispersity (Table 7.3(b)).

Furthermore, other thermal properties also vary with DP and polydispersity. Using cationic polymerization, Sagane and Lenz37 prepared poly(vinyl ether) combs having methoxybiphenyl side groups (cf. Section 7.5.2.1). Not only were mesophase identities and transition temperatures variable but so also were the associated enthalpy (IlH) and entropy (IlS) changes as seen in Table 7.4.

Finally, Haase et al. 38 have demonstrated a correlation between DP and dielectric phenomena. Table 7.5 contains data obtained using cyanobiphenyl acrylic combs for the energy of activation (EA ) of dielectric relaxation at 365 K from a state of alignment parallel with an electric field. An inverse relationship between EA and DP is indicated; the transition temperatures tend to increase with DP as noted previously.

7.3.1.3 Tacticity One might expect tacticity, the three-dimensional relationship between bond dispositions at the points of attachment onto the backbone, to influence the ordering of mesogens (especially with short spacers). Relatively little has been published about this, but Pugh and Percec lO have noted that tacticity seems to affect mainly crystalline melting points, which increase with stereoregularity, thereby lowering the


Table 7.3 Influence of DP and polydispersity on phase

transition temperatures (K) of

Ie Me3SiO-f-Si-O+SiMe3

I (CH 2)6@COO@CN

(a) DP

VP Mw/Mn Tg Tel

30 1'15 275 401 55 1·15 282 410 84 1·06 286 415

107 1·12 287 418

(b) Polydispersity

Mw/Mn Tg Tel ~--- ------~ ---

15·5 261 390 1·9 270 396 1·15 275 401

Table 7.4 Comparative thermal data for two specimens of

+ CH2-CH+ bCH2CH20@--@-OMe

363

Transition temp. (K) AH(J g-l) AS x 102(1 g-l K- 1 )

Mn x 10- 3 Mw/Mn Ts-N TN-i AHs-N AHN-i ASs-N ASN_i -------~-

6·7 2·5 458 11·9 2'6 4·3 1·2 418 433 H 4·6 1·8 1-1

meso phase range (Tel remains more or less constant). For glassy materials Finkelmann has proposed a similar relationship between Tg and TcJ with increasing stereoregularity.31

Usually free-radical polymerization of acrylic systems produces mainly stereo random (atactic or heterotactic) polymers, but Duran and Gramain have demonstrated a correlation between tacticity triads and the spacer length in acrylate monomers subjected to free-radical solution

364

fjp

14 21

150

Derek J. Simmonds

Table 7.5 Influence of OP on thermal and dielectric parameters of

+CH2-CH+

bOO<CH2l60--@--@-CN

Phase transition temperatures (K)

g 305 RN 327 SA 368 N 383 i g 305 RN 327 SA 371 N 384 i g 305 RN 353 SA 398 N 405 i

144 140 126

polymerization (Fig. 7.6).39 The authors suggest that stereoregular dispositions (isotactic and syndiotactic) are influenced by biphenyl-biphenyl interactions during polymerization; this may be a fairly general phenomenon in the synthesis of comb LCPs from LC monomeric reactants.

Me

ACOO-<CH2CH20ln-<Q)-<Q)-OMe

SYNDIO

40 ~

-! HETERO . ~ A .. 1/1 " * "0 A

" 30 'i: I-

3 4 Homopolymer spacer length (n)

Fig. 7.6. Effect of spacer length on tacticity in the free-radical polymerization of a methacrylate (After Ref. 39).


In new synthetic work, Mallon and Kantor have used coordination polymerization to prepare stereoregular hydrocarbon combs (65-90% isotactic) from monomers that are 4,4' -dialkylbiphenyls containing a terminal double bond in the longer alkyl chain.40 The resulting polymers exhibit smectic Band smectic E meso phases that are uncommon in comb LCPS. 14a

7.3.2 The Flexible Spacer General comments have already been made (Section 7.2.2) and we noted that the linkages, from backbone to flexible spacer and then onto mesogen, exert influences on transition temperatures. Figure 7.7 illustrates effects on both Tg and Tel when an ether connection is replaced by a methylene group (direct, alkyl connection).41 The constitution of the spacer chain affects LCP properties also,l 0 but a particularly influential variable, once a spacer-linkage combination has been selected, is the length of the spacer chain. Here we shall concentrate on that aspect, and then look at the mode of attachment onto the mesogen lath since there has been recent interest in lateral attachment (producing 'parallel' polymers-Table 1.1).

Fig. 7.7. Effect of changing the mesogen linkage.

7.3.2.1 Spacer length In a series of comb LCPs where spacer length is the only variable, there may be two reasons for any observed differences. Firstly, we may be seeing the result of changing the effective length of the mesogen (i.e. at least a part of the spacer is involved in the mesogenic structure) which is known to affect both the nature and thermal stability of mesophases (Section 7.3.3.1). Secondly, we may be seeing


the result of differing degrees of freedom of mesogen motion with respect to backbone motion, i.e. decoupling via the spacer effect. In any case the choice of spacer length is one of the most effective determinants for the synthetic chemist to exploit in optimizing thermal properties.

A comprehensive programme of synthesis and characterization by the Bordeaux group42 has provided a wealth of comparative data on which Table 7.6 is based; note that the polymers are well defined, monodisperse systems. The general points to emerge from these and comparable data are as follows:

- As spacer length increases so Tg tends downwards. - An odd-even effect (values of n) often underlines general trends in

comparable transition temperatures (e.g. X=OMe; Tel for n=3, 4, 5, 6); see below for another example.

- Any nematic phases normally arise at shorter spacer lengths, longer spacers tend to encourage smectic phases.

Table 7.6 Effect of spacer length on thermal properties of siloxane combs

Me

I Me,S;O+r -olS;Me, ° (CH2)40-@-O-~-©-X

n Phase transition temperatures (K) !1HcM g-l)

X=CN 3 g 309 K 373 SA 449 i 4·6 4 g313 K351 SA 447 I 5·0 5 g 297 K 377 SA 457 i 5·6 6 g 300 K 355 SA 463 i 6·0 8 g298 _~ SA463i 8·5

10 g291 K381 SA468i' 7 11 g 290 K 355 Sc 392 SA 474 i 9·2

X=OMe 3 g 288 SA 396i 2·0 4 g 280 SA 347 N 377 i 1·6 5 g 277 K 344 SA 395 i lO 6 g 269 SA 383 i H 8 g268 K 318 SA 400i 4·5

10 g268 K 315 SA 406i 5·5 11 g 297 K 239 Sc 333 SA 407 i 6·1


Additional smectic phases may appear as spacer length mcreases (e.g. Sc at n = 11).

- Where LCPs are potentially crystalline, short spacers are required to suppress crystallinity (e.g. X = OMe; n = 3, 4).

- Clearing enthalpies tend to increase with spacer length.

The data also illustrate the polymer effect, viz. tethering a mesogen to a polymer tends to increase the meso phase ordering. Thus, of the smallmolecule species (Fig. 7.8) used in the synthesis of the smectic combs (Table 7.6), all are nematic apart from (X = CN, n = 11), and only the cyanomesogens with long tails (n > 8) exhibit smectic as well as nematic phases.

Fig. 7.S. Monomers of the combs shown in Table 7.6.

Similar trends have been noted in many other series involving different polymer components. For example, Fig. 7.9 presents data on a homologous sequence of methacrylates where the mesogen is a cholesteryl ester,43 and clearly illustrates both the gradual decrease in Tg and an odd-even effect on Tel. Cholesteryl esters as SMLCs are associated with cholesteric phases (chiral nematic ordering), but the polymers are smectic (with an additional mesophase appearing at n = 9), as we might predict on the basis of a polymer effect. Once again we might look for variations in other relevant parameters as well as thermal transitions. The polymers described by Table 7.744 exhibit the anticipated trend in Tg and in the attainment of greater ordering with longer spacers. In addition, the activation energies (E A ) of dielectric relaxation from homeotropic alignment with an a.c. field appear to vary with spacer length. Simon and Coles44 associate their observation with increasing freedom of the side groups with longer spacers; a significant drop in EA is apparent between the cases n = 6 and n = 8.

7.3.2.2 Lateral attachment of mesogens: parallel (biparallen systems While we have not otherwise looked at connections between spacer and mesogen, there has been particular interest in LCPs with laterally attached meso gens since Hessel and Finkelmann reported the first example in 1984.45 Table 7.8 summarizes published data on several polymers that have been prepared since then. In general the phase

368

u . ~

250

200

~ 150 J .... IV L 41 a.

~ 100 ~

50

Derek J. Simmonds

Q , , d 'p

A \

b" \ \ \

0'

\ , a

Isotropic

'0.0'0- '0

o \

'. Mesophase

'. Mesophase \

b lJ..,

4 , , \

~,

/S'n_ -a --A-~ 15. -- - --6.

Glass

Carbon number (n)

Fig. 7.9. Homologous cholesteryl methacrylate combs (After Ref. 43).

Table 7.7 Thermal and dielectric relaxation data for siloxane combs

Me

Me,s;of t -ofS;Me,

I ~ W --<n\-(CH2 )nO frc- o ~ CN

Me

n Phase transition temperatures (K) Low temp. Temp. = 0·98 Tel

4 g297i 5 g 292 N 338 i 202 172 6 g280 N 321 i 238 190 7 g282 S 360i 8 g282 S 375 i 146 95

11 g 281 K 318 S 404 i

n m

0 1 1

1 2

1 4

4 4

4 6

6 4

10

4 10

6

11

1 11

2

11

3 11

4

11

5 11

6

11

7 11

8

Tab

le 7

.8

Pha

se t

rans

itio

n te

mpe

ratu

res

(K)

of s

ome

para

llel

LC

Ps

wit

h la

tera

l m

esog

en a

ttac

hm

ent

R =

Cm

H2

m+ 1

0-@-coo~ooc-@-ocmH2m+1

~ H

I ~e

Cl I

Me

I

-+-C

H2-

C+

I

+C

H2-C

+

I +

CH

2-C

-+

-

I -+

CH2

-C

+

I -+

-5

;-0

+

I R

C

OO

(CH

2lnR

COO(CH~nR

CO

O(C

H2l

nR

(CH

2ln

OC

OR

Ref

eren

ce

g 39

8 N

600

Dec

."

46

g 37

6 N

433

i

g 37

1 N

438

i

g3

70

N4

36

i 47

g 29

2 N

322

i (g

285

N 3

47)

K 3

64 ib

g2

84

N3

31

i g

277

N 3

40 i

g 26

4 N

332

i

g 26

0 K

306

N 3

34 i

g 30

8 N

392

i 34

g

29

6 N

372

i

g3

13

N3

40

i 32

(g 3

07 N

355

) K

365

iC

g 29

9 N

331

i

g2

92

N 3

40

i g

283

N 3

24 i

g 27

8 N

333

i

g 27

4 N

321

i

g2

71

N 3

34

i

g2

90

N3

61

i 34

g

28

6 N

36

4i

(g 3

04 N

352

) K

365

iC

g 29

3 N

339

i 3

2,3

3

g 27

3 N

333

i

"Pre

lim

inar

y da

ta,

to b

e co

nfir

med

. bL

C p

hase

det

ecte

d on

ly o

n c

ooli

ng d

ue t

o ra

pid

crys

tall

izat

ion.

C

Part

ial

crys

tall

izat

ion

afte

r an

neal

ing

for

24 h

.

370 Derek 1. Simmonds

behaviour of meso gens is much less affected by lateral attachment to a polymer, than by the usual longitudinal attachment where we have already noted significant polymer effects.

Three aspects of the phase behaviour of these polymers are especially noteworthy. Firstly, there is no apparent tendency towards smectic behaviour, even with long spacers. Long mesogen tails (large values of m), which confer smectic behaviour in SMLCs and conventional comb LCPs, are not smectogenic here. Indeed, no smectic parallel LCPs have yet been reported despite deliberate incorporation of normally smectogenic features (e.g. polar mesogen tails, copolymerization with methyl methacrylate, etc.).32 Clearly, lateral substitution is incompatible with the highly ordered packing required in the formation of smectic phases.

Secondly, we see in Table 7.8 a rare example of a side chain polymer with no spacer that is liquid crystalline (n=O, m= 1). Zhou48 uses the term 'mesogen-jacketed' to describe the probable microenvironment around the backbone when spacer lengths to lateral mesogens are short, with close proximity of the centres of gravity of the mesogens and the points of substitution onto the backbone. The polymers are then somewhat reminiscent of the longitudinal ('worm') LCPs dealt with in Chapter 8.

Finally, attachment of spacer/linking groups onto the side of the mesogens, rather than at their ends, inhibits their rotational freedom. The ordering of a conventional (uniaxial) nematic phase involves alignment with respect to the mesogen long axes only; rotational freedom prevents biaxial alignment where there is also ordering with respect to the short axes. Careful analysis of mesophase textures32-34 suggests that some of the polymers in Table 7.8 afford biaxial nematic phases. Keller et al. have identified nematic-to-nematic transitions among the siloxane polymers (n = 4) by monitoring macroscopic magnetic susceptibility (AX) versus temperature in the nematic region (optical texture) as shown in Fig. 7.10. 34

7.3.3 The Mesogen We have dealt somewhat artificially with flexible spacers as entItIes separate from the mesogens, although we noted (Section 7.3.2.1) their influence in defining the overall length of mesogens. Certainly mesogen length is a significant parameter, but we must also consider the effects of, for example, mesogen flexibility and dipole moment, which will be affected by the nature of linking and substituent groups within the

E ,·0 GJ

f'. o

~ 0'5


Fig. 7.10. Nematic to nematic transition (N -> N').

mesogenic structure. To some extent structure-property trends accepted for SMLCs49 hold good for comb LCPs; the major issues have been thoroughly reviewed elsewhere10•11 and a brief resume will suffice here.

7.3.3.1 Mesogen length Increasing the mesogen length leads to effects already discussed with respect to spacer length (Section 7.3.2.1). Lengthening the terminal tail (e.g. an alkyl group at the end remote from the spacer group) has a similar effect, and analogous trends are evident in lateral systems such as the polymers shown in Table 7.S. Thus Fig. 7.11, reproduced from Ref. 33, illustrates the gradual lowering of Tg and the oddeven effect upon Tel as m varies in the homologous series of methacrylate polymers (n= 11, m=2-S) shown in Table 7.S. Examples of polymers with increasing mesogen core lengths are shown in Fig. 7.12,50 and clearing temperatures show a steady rise with n. For example, when X = 0 Me, Tel is reported as 334, 592 and 633 K for n = 0, 1 and 2, respectively.

7.3.3.2 Terminal substituent Calami tic mesogens are generally lath-shaped with the spacer group attached at one end of the rigid lath. The identity of the terminal group, attached at the other end of the lath, is an important determinant of both the magnitude and direction of the mesogen's dipole moment. Thus, both self-organization and field-induced alignment of mesogens

372

350

340

330

320

~ 310 .....

300

290

280

270

Derek J. Simmonds

Fig. 7.11. Tg and Tel variation with m; methacryaltes of Table 7.8 (After Ref. 33).

Me

I +Si-O-+-

o

(C~)30-@-g-O~x

Fig. 7.12. Variation of the rigid core length.

are influenced by the choice of terminal substituent, with attendant implications for phase-transition temperatures and realignment phenomena.

The effects on transition temperatures of four common terminal groups are shown in Table 7.9, which illustrates three of the following four general trends:42

- The less polar substituents (alkyl, alkoxy) favour nematic phases (whose appearance also depends on spacer length).

- Cyano and nitro mesogens are smectogenic and are often associated with similar clearing temperatures (but different clearing enthalpies AH cd in similar polymers.

- Clearing temperature increases with terminal group polarity.


Table 7.9 Effect of the terminal substituent on thermal properties

Me

I Me3SiO +Si-OtsiMe3

I 35 ° (CH2)nO-@-O-g-@-x X Phase transition temperatures (K) "'HeM g-l)

eN g 313 K 351 SA 447 i 5·0 N02 g 293 SA 438 i 2·7 OMe g 280 SA 347 N 377 i 1·6 Me g 277 N 332 i 2·0

Unsubstituted systems (X = H, not shown in Table 7.9) are not liq uid-crystalline.

These data concern substituent polarity. Another important variablechirality-is not covered here; we have mentioned chiral mesophases already (Section 7.2.3) and will return to chiral smectic C materials later.

7.3.3.3 Lateral substituent The inclusion of a lateral substituent on the mesogen rings can have a marked effect and could be a useful device in the tailoring of mesogen design to selected properties. An important phenomenon was reported by Gemmel et al.;41 Table 7.10 presents some of the observations. Taking the entries in pairs, we see that a lateral methyl group significantly lowers the degree of ordering exhibited by a polymer. The substituent suppresses mesophase formation at low spacer length, suppresses crystalline and smectic properties to give glassy nematics when n = 5, 6, and suppresses crystallinity even beyond that (n = 8). The suppression of crystallinity was ascribed41

to the inhibition of crystalline packing by the methyl substituent, presumably a steric effect. Electronic effects may be involved in the smectic/nematic variations, as has been suggested for the malonate combs shown in Table 7.11, where a lateral dipole is introduced without significant steric variation by the use of a pyridine ring in place of a benzene ring. 51 The strong lateral dipole of the 2,5-disubstituted pyridine is invoked to explain the smectogenic effect of the

374

n

4 4 5 5 6 6 8 8

11 11

Derek J. Simmonds

Table 7.10 Effect of a lateral methyl substituent

r MO,S;O+ r -olS;Mo, °

(CH2)nO-@-~-O?-CN R

R Phase transition temperatures (K)

H g 299 SA 423 i Me Tg=297 H K 380 SA 456·5 i Me g 292 N 338 i H K 328 SA 454·5 i Me g 280 N 321 i H K 335 SA 467 i Me g 282 SA 375 i H K 349 SA 471·5 i Me g 281 K 318 SA 404 i

Table 7.11 Benzene/pyridine ring: lateral effects

o 0 II II

+C-CH-C-O-Z -0+

I (CH2)sO-@-N=CH-@-N02

X

z X Clearing transition temperature (K)

CH2CH2CH2CH2 CH N 358·9 i SA 315·2 i N* 326·8 i S~ 347-2 i N* 308·9 i S~ 346·2 i

CH2CH2CH2CH2 N CH2CH2CH(CH3) CH CH2CH2CH(CH3) N CH2CH2CH(CH3)CH2CH2CH2 CH CH2CH2CH(CH3)CH2CH2CHz N

pyridine systems. Starred phases are chiral owing to the use of optically active diols in some polycondensations. No molecular masses were quoted, but all the polymers shown had received similar, prolonged heat treatment.


7.3.3.4 Central linkage Two series of siloxane polymers, corresponding to two spacer lengths, are shown in Table 7.12, where four central mesogen linkages can be compared. Absolute comparisons are subject to error since for the biphenyl systems41 x = 50 while the ester and ether examples42

have x = 35. However, the final parenthetical biphenyl entry,52 for which x = 36, suggests that clearing points at least should be closely comparable. While the data concern only cyano-terminated mesogens, some general points are indicated; in particular we can conclude that the central linkage is indeed a determinant of thermal behaviour. In

n Z

0 II

4 O-C 0 II

4 C-O 4 -

4 OCHz 0 II

6 C-O 0 II

6 O-C 6 -

6 OCHz (4 -

Table 7.12 Effect of the central linkage Me

M""O+fi -otsiMe3 <CH2lnO--@-Z--@-CN

Transition temperatures (K) Reference --------

g 313 K 351 SA 447 i 42

g288SA412i 42

g 301 SA 405·5 i 41

g 283 K 351 i 42

g 300 K 355 SA 463 i 42

g 289 K 317 SA 443 i 42

g 286·5 SA 448-5 i 41

g273K318SA359i 42

g287 SA 405 i 52)


these examples it seems that clearing temperatures drop in the order

o 0

11>11 > O-C C-O '" biphenyl OCH 2

The most flexible, the benzyl ether linkage (OCH 2 ), is here associated with the lowest Tg and Tel values and fails to provide a mesophase when n = 4. Only the cyanobiphenyls are non-crystalline at both spacer lengths considered, while all four benzyl ethers and cyanobenzoate esters exhibit crystalline phases.

The inclusion of flexible linking groups within the mesogen implies the possibility of conformational isomerism that may suppress crystallization of the meso genic side groups, although backbone crystallization could still occur (the crystalline structures in Table 7.12, Z =

OCH 2 , may be of this form). It is important then to consider whether comb LCPs are capable of biphasic behaviour, that is independent phase behaviour of either the backbone or the substituents. Percec53

has pointed out that virtually complete decoupling of backbone and mesogen, via long flexible spacers, could lead to just such a phase separation if there is mutual immiscibility of the two structural components. In such cases new conformationally flexible components of mesogens could help avoid unwanted mesogen crystallization when long spacers are required for property optimization. In Fig. 7.13 three such mesogens are shown that contain units capable of conformational isomerism, namely diarylethanes (1)54, 2-aryl-l,3-dioxo-2-

Fig. 7.13. Mesogens capable of conformational isomerism.


borinanes (11)55 and 2,5-diaryl-l,3-dioxanes (IIW 5, which are also available as cis/trans geometric isomers.

7.4 COPOLYMERS

In small-molecule LC systems, mixtures containing several LC components tend to be more useful for commercial applications than are pure LC substances. 8 By analogy, composite polymers may, despite expense, turn out to be the most useful comb systems. In such composites the mixing may be purely physical (e.g. blends of comb LCPs with dyestuffs, as mentioned in Section 7.6.1), or the components may be chemically bonded into copolymers. In so far as copolymerization of two monomers generally leads to properties that are modified (but not fundamentally different) with respect to either of the analogous homopolymers, we may expect that finetuning of comb LCP properties should be feasible by copolymerization (e.g. by utilizing the sort of structure-property correlation data we have just derived for homopolymers).

Two basic designs of dual copolymer comb LCP can be envisaged; combining mesogenic units with non-mesogenic units, or combining two different mesogenic units. A special case is the copolymerization of mesogenic monomers with appropriate cross-linking monomers to generate network materials (cf. Section 7.4.3).

Copolymerization raises serious issues concerning structure. Not only are the homopolymer variables still relevant (DP, tacticity, etc.), but we must add copolymer composition and the order of substitution along the copolymer chain to the list. Even so, sufficient is now known to allow a brief review of copolymer properties.

7.4.1 Copolymers of Mesogens with Non-mesogens Although the dilution of mesogens, via the incorporation of nonmeso genic units, generally has only minor effects on the nature of the mesophases observed, the changes in phase-transition temperatures can be quite marked depending on the ratio of mesogenic to non-mesogenic units. For example, Table 7.13 shows a series of copolysiloxanes 56 where the alkylmalonate diester mesogens are diluted by the inclusion of dimethylsiloxane units in the chains (these are 'double combs' according to the Brostow classification of Chapter 1). Copolymerization results in the loss of a short-lived nematic phase


Table 7.13 Siloxane copolymers with malonate diester meso gens

Me Me I I

Me3SiO-H-Si-O+,---+Si-O+YlxSiMe3

I Me

?i ....IR'L ~ ....IR'L (CH2)3CH(C-O-CH2CH2 Qr--c-o Qr--0CH3)2

y Transition temperatures (K) I'>.T

0 g 306 SA 364 N 371 i 1 K 303 SA 363 i 60 5 K 321 SA 387 i 66

10 K303SA380i 67

and induces crystallinity, but the copolymers show consistent phase behaviour, albeit with varying temperature profiles. In this and some other copolysiloxane series involving dimethylsiloxane diluents, the longevity of the mesophase (L\ T) appears to increase as y increases (up to 10 at least). However, in other studies,S? the general trend has been for Tel to decrease as the diluent content increases.

The size of the diluent side group is reflected in the extent to which comparable transition temperatures are affected. For example, in a series of methacrylate copolymers involving cholesteryl ester meso gens with alkyl ester diluents (Fig. 7.14),58 the magnitude of the depression of Tel (at constant molar fraction of diluent) increased in the order R = methyl < n-butyl < t-butyl < dodecyl < octadecyl.

Me Me I I

-f-CH2-C-+------f-CH2-C+ I I

COOR COO(CH2)1O-COO-Cholesteryl

Fig. 7.14. Cholesteryl ester copolymers.

Dilution, then, can be useful in tailoring a mesophase profile to a desired temperature range, perhaps extending the range of applicable mesogens. New structure-property correlations have been delineated by the Hull group59 using consistent and well-characterized copolysiloxanes (with dimethylsiloxane diluents) to compare new fluoromesogens with cyano analogues. Figure 7.15 is drawn from their work and shows what can be achieved, in this case a detailed


Me Me I I

Me3SiO+ Si -0-r,. ---- +Si -0 +'y SiMe3 I I

Mesogen Me

U 120 !, 1/1 100 TSA_i ~ ::l 80

Ref ... " L. 41 60 Q. 70· E ~ 40 c .Q 20 33·

7· ... 'iii Tm Ref c 0 5· " Tg Ref L. I- -20

X: CF3

Fig. 7.15. Mesogen phase correlations using well-characterized siloxanes (After Ref. 59).

analysis of structure correlations between the fluoromesogens and a reference, cyanoester system (represented by the horizontals) undergoing g-K (Tg), K-SA (Tm) and SA-i (Tel) transitions.

Mesogen dilution can also affect the miscibility of mesogenic components with backbone components of comb LCPs. Biphasic behaviour (i.e. phase separation) has been noted in comb polymers60 and recent work suggests that adjustment of mesogen/diluent ratio can influence this aspect of comb morphology,61 once again offering a fine-tuning mechanism (cf. Section 7.6.1.2).

7.4.2 Copolymers with Two Different Mesogens Structural complexity abounds with this design since any two meso gens can, in principle, be combined and the two meso gens can be associated with different spacers and/or tail groups. There can, accordingly, be an


element of serendipity in the choice of comonomer combination, but the guiding principle seems to hold good of using a second mesogen whose homopolymer has some desired property to modify the behaviour of a polymer of interest. Thus, for example, acrylic homopolymers containing cholesteryl ester meso gens tend to be smectic but the inclusion of nematogenic p-methoxyphenyl benzoate units into copolymers generates cholesteric, rather than smectic, materials. 62

Recently Finkelmann et al. 63 reported that a similar comesogen pairing, but based on a siloxane rather than acrylic backbone, generates a blue phase64 after thorough annealing; this is the first blue phase reported for a comb LCP.

We have already remarked that copolymers may come to dominate the applications of comb LCPs (cf. Section 7.6). Once again the Hull group, using well-characterized systems, have demonstrated some possibilities. Cyanomesogens are smectogenic (tending to give rather viscous meso phases) but provide a useful, strong positive L1e (desirable for efficient mesophase switching). In separate studies, both briefly reported in Ref. 59, a cyanobiphenyl copolymer with comesogens containing lateral methyl substituents (Section 7.3.3.5) and a cyanophenyl benzoate copolymer with laterally attached (i.e. 'parallel') comesogens (Section 7.3.2.2) both provided nematic materials with potentially high positive L1e.

A series of elegant studies by Percec et al. has demonstrated some subtle, but significant, effects in copolymers based upon constitutional isomer pairs in methacrylate, acrylate and siloxane materials. The acrylic copolymers tended to be non-crystalline but the more flexible siloxanes were all crystalline, whatever the spacer length (Fig. 7.16, x # 0, Y # 0, n> 3).65 Controlled synthesis provided three series of materials (n = 3, 6 and 11) each including both homopolymers (x = 0, y = 100;66 and x = 100, Y = 067 ) and copolymers of known composition.68 Each series provided its own information as follows.

The n = 3 series was nematic and demonstrated that copolymerization depressed both K-N and N-i transition temperatures to similar extents (little effect on the overall persistence of the mesophase). However, the 11 = 11 series was smectic (as expected for long spacers) and exhibited a dramatic improvement in stability of the smectic phase. The S-i temperatures were little changed from the homopolymer values (411 K and 416 K for x=O and y=O, respectively), while the K-S temperatures were greatly depressed (400 K for x = 0 and for y = 0, but about 360 K for x = y = 50). The n = 6 series was (like n = 3) nematic, but (like n = 11) demonstrated


Me Me I I

Me3SiO+Si-O+.-----Si-O+.SiMe I x I Y 3

(CH 2)n (CH 2)n

I I ° 0

x+y:100

n : 3,6,11

Fig. 7.16. Homo- and copolysiloxanes based on constitutional isomers.

381

a marked depression of crystal-to-mesophase transition temperature, much greater than the effect on Tel; a significant increase in meso phase stability resulting from copolymerization. The result was that while both homopolymers (i.e. for n = 6) were monotropic, exhibiting a nematic phase only during cooling cycles, the 50/50 copolymer was enantiotropic. The relevant temperature gaps (.1 T) were 24 K (x = 0), 22 K (y = 0) but 44 K (x = y = 50), as noted during cooling experiments.

7.4.3 Cross-linked Polymers-Network LC Elastomers If, instead of copolymerizing two different mesogenic monomers, one mesogenic monomer is polymerized in the presence of a bifunctional (but non-mesogenic) monomer of similar reactivity, then a cross-linked network polymer will result (Fig. 7.17), the degree of cross-linking reflecting molar proportions of meso genic and cross-linking monomers used. The first network LCP was a siloxane elastomer reported by Finkelmann et al.,69 and since then network acrylates, 70. 71 methacrylates 71 and malonates 72 have also been prepared. Recently chiral LC malonate elastomers showing chiral smectic and nematic mesophases were announced. 73 The field has been thoroughly reviewed by Gleim and Finkelmann 74 and by Zen tel. 7 5

Moderate cross-linking ("" 10-20%) has no effect on the nature of the mesophases compared to an analogous homopolymer, but tends to depress all phase transition temperatures by a few degrees. Of greater significance are the mechanical properties and related alignment phenomena. Above Tel acrylic network LCPs exhibit normal rubber elasticity,


~c=O +~C=O I I

~--------~

~ 1 C=O C=O I I o 0

Me I

I o=c~

+Si-O+ + I H

(a)

+

Me Me I I

o 0 I I

C=O C=O

~ .... ~

+Si-O+------------+Si-O+

+ Si-O+---+Si-O+ I I

Me Me

(b)

Fig. 7.17. Schematic representation of acrylic (a) and siloxane (b) network LCs.


but anomalies are observed at mesophase temperatures--e.g. high levels of mesogen orientation at relatively low strain. 76 Similar thermoelastic (as well as photoelastic) parallels with conventional elastomers have been observed with siloxane network LCPs. 77 The detailed elastic response to mechanical deformation depends on the detailed structure of the network LCp. 74 In particular, a dependence on the spacer length (backbone to mesogen) of stress-optical and thermomechanical measurements has been noted in the case of siloxane systems. 78 Depending on, for example, the spacer length, orientation may be either parallel or perpendicular to the direction of the mechanical force, and stretching and compression tend to have opposite effects on alignment direction.

The optical and elastic properties of these novel network LCs are quite new. Orientation via mechanical deformation, and the resultant stressoptical phenomena, can be added to the range of behaviours already associated with side chain mesogenic polymers, such as electro-optic and magneto-optic phenomena. The comb LCP design is clearly capable of delivering a wide range of novel materials.

7.5 SYNTHESIS OF COMB AND PARALLEL SYSTEMS

The principal synthetic approaches are illustrated in Fig. 7.18. A side chain polymer can be made by the polymerization of an appropriately substituted monomer using addition (approach a) or condensation polymerization (approach b, where X and Yare mutually reactive functional groups). Alternatively, the compounds can be prepared via the attachment of the desired mesogenic entities (generally including a spacer group) onto a preformed polymer in polymer homologous reactions (approach c) that again require mutual functionality. All three approaches have been the subject of comprehensive, recent reviews covering both hydrocarbon-based polymers (approach a, b or C)10 and siloxane polymers (c),u

The construction of the mesogenic entities is in the realm of organic synthesis, while preformed polymers are often commercially available or are prepared using established polymer techniques; these aspects are not, on the whole, covered here. It is, however, important to emphasize that all components used should be rigorously purified and, in the case of preformed polymers, fractionation to a known molecular mass range is desirable. The incipient mesogens need not always give mesophases prior to comb formation; they will be subject to the polymer effect and show


+

y

~ y

Fig. 7.18. Synthetic approaches to comb LCPs.

an increased tendency towards ordered-phase behaviour as a result of polymer formation.

7.5.1 Addition Polymerization of Mesogenic Monomers Addition polymerization of alkenes can be performed in several ways. Most commonly polymerization is initiated using reactive free radicals provided by an added reagent. Initiation is usually achieved thermally, but photochemical initiation is also relevant here. Initiation procedures using anionic or cationic reagents are attractive alternatives since they offer better control of molecular mass; some cationic procedures are mentioned below (Section 7.5.2.1).

7.5.1.1 Thermal free-radical polymerization Figure 7.19 illustrates a typical procedure for the synthesis of a methacrylate LCP. The monomer can be prepared, as shown, by esterification of methacrylyl chloride. Polymerizations are usually initiated via thermal homolysis of azobisisobutyronitrile (AIBN) in a solvent such as toluene, chlorobenzene or tetrahydrofuran. After termination, the polymer is isolated using several reprecipitations from a non-solvent such as methanol. Techniques such as gel-permeation chromatography, high-performance liquid chromatography or thin-layer chromatography should be used


AIBN/60°C •

Toluene

Fig. 7.19. Synthesis of a methacrylate comb.

385

to establish purity and determine the molecular mass and polydispersity. Products are commonly annealed at a temperature a little below their clearing temperature.

The method can be adapted to the preparation of copolymers or cross-linked network polymers. For copolymerization it is desirable to use monomers of very similar reactivity so that their known molar proportion will be accurately expressed in the copolymer composition; this should also be checked, for example using spectroscopy. So long as monomers are of similar structure near the reactive double bond (e.g. both methacrylates with several CH 2 groups in the spacer) then their reactivities should be virtually identical.

For network polymer formation a small amount of an appropriate divinyl monomer can be included in the monomer feed as a cross-linking agent. By analogy with copolymerization, as long as the monomer and cross-linker are of similar reactivity then the degree of cross-linking will be statistical and can be controlled by the molar proportion of monomer to cross-linker. An example of the preparation of acrylic network LCPs is shown in Fig. 7.20. 79

7.5.1.2 Photochemical radical polymerization Polymerization can be initiated using radicals generated via the UV irradiation of appropriate initiators in the presence of the monomer. This approach to comb LCPs, first reported by Broer et al.,80 is attractive since it can be performed at a preselected temperature. Therefore, monomers can be polymerized in an aligned LC phase and the effects of ordered polymerization can be examined.

Bulk polymerization of acrylate (I) shown in Fig. 7.21 was carried out8 !

using either of the initiators (II) or (III). The monomer (I) exhibited monotropic nematic and smectic phases on cooling from the isotropic

386

+

Derek J. Simmonds

AIBN 155°C. Elastomer chlorobenzene

Fig. 7.20. Synthesis of a network LCP.

____ ~h~v __________ -+. combLCP

o R1

o-g-i:-R2

~ ~1 (II) R1: Me, R2: OH

(III) R1: OMe, R2: Ph

Fig. 7.21. Photopolymerization.

liquid phase, and a stable nematic phase when heated from the crystalline state. Photopolymerization from either of the monotropic phases was hindered by crystallization of the metastable LC phase, leading to incomplete conversion (60-90%). Photopolymerization from the stable nematic phase was rapid and complete, leading to polymer in a smectic phase (polymer effect). Polymerization from the isotropic phase went to completion but was much slower than either of the former procedures. Therefore, monomer ordering is a significant factor in comb LCP


synthesis. Photoinitiation offered a very rapid polymerization, e.g. 15 seconds to 1 minute in this study.

7.5.2 Ionic Polymerization Initiation by ionic species often produces well-controlled polymerization that is not subject to spontaneous termination or chain transfer. It is sometimes possible to produce predictable and narrow molecular mass ranges by incorporating stoichiometric quantities of end-capping reagents to encourage termination at a desired degree of polymerization (DP) assuming complete consumption of monomer. Ionic initiation can also produce 'living' polymers that are still reactive (unterminated) after full consumption of monomer. The DP at that point is known from the ratio of ionic initiator to monomer, so that living polymerization is ideal for producing narrow molecular mass ranges and can also produce block copolymers by the chain growth of different monomers added in sequence to the living system. Unfortunately, anionic initiators are generally too basic and too nucleophilic to use with monomers of interest in LCP synthesis; side reactions involving the mesogens tend to compete with the desired initiation reaction.

7.5.2.1 Cationic polymerization Cationic polymerization of comb LCPs was first reported by Percec82

who used the Lewis acid boron trifluoride etherate as initiator. This approach, illustrated in Fig. 7.22, gave useful solution preparations of polyvinylethers (R = H) and polypropenylethers (R = CH3)' Recently Sagane and Lenz37 demonstrated that the use of the cationic initiator system HI/I2 brought about living polymerizations of similar monomers and greatly improved the molecular mass distribution of the products.

Percec has extended the use of cationic procedures to the synthesis of vinyl ether copolymers83 based on constitutional isomers, analogous to the copolysiloxanes shown in Fig. 7.16 (Section 7.4.2). Other work, by Percec and Hahn, will prove useful in the controlled synthesis of siloxane backbones (Fig. 7.23).36 Cationic, ring-opening polymerization of cyclic tetrameric siloxanes, in the presence of an end capper (hexamethyldisiloxane) and initiated using triflic acid, can be used to prepare homo- (Y = n = 0) or copolysiloxanes whose DP and compositions are determined by the ratio of end-capper to monomer in the feedstock. These polymers are starting materials for siloxane comb LCPs (Section 7.5.5).

388

OMe

Derek J. Simmonds

cationic

initiator

~ +CH--CH+

I

..

o I

(CH 2)n I o

OMe

Fig. 7.22. Cationic polymerization of meso genic polyvinylethers.

~ X \.~i-O< + y

I H

Me Me I I

~ \.~i-O~

I Me

Me3SiO-+ti-O+ni---+ ti - O -+"SiMe3

H Me

Fig. 7.23. Cationic polymerization of siloxanes.

7.5.3 Condensation Polymerization Condensation polymerization is of greatest utility for main chain mesogenic systems (longitudinal polymers in Table 1.1) or combined main and side chain mesogenic systems (double polymers in Table 1.1), where polyester or polyamide backbones are common. Zentel72 has extended the method to produce cross-linkable combined polymers (e.g. see Fig. 7.24), but other polymalonates, with side chain mesogens only, have also been prepared in this way.84 Combined (longitudinal) as well as comb LC polyurethanes have been prepared via the condensation polymerization of mesogenic or non-mesogenic diisocyanates with diols bearing mesogenic side groups.18


+

I

Fig. 7.24. Synthesis of a cross-linkable combined polymer.

7.5.4 Polymer Modification Reactions The modification of preformed polymers via the attachment of meso genic entities (polymer homologous reactions) is a useful alternative to the chain polymerization of vinyl systems and an essential strategy for siloxane combs. However, careful monitoring is necessary to ensure complete reaction, and any unreacted mesogen (often used in excess) must be rigorously removed from the product. We can illustrate the approach for phosphazene combs as in Fig. 7.25,17 where mesogenjspacer species containing a terminal sodium alkoxide functionality (RONa) are shown displacing chloride ions from either cyclic trimers or high polymers to give mesomorphic products.

Solvent selection can be a problem in these reactions owing to the dissimilarity of the reactants, usually a polar nucleophile and a non-polar substrate. Keller85 has applied phase-transfer catalysis (PTC) to polymer homologous reactions, and has demonstrated its utility in the preparation of comb LCPs based on polyacrylate, polymethacrylate, polyitaconate, polymaleate and poly(methylvinylether-co-maleate) backbones86 via ionization of the poly(carboxylic acid) compounds and esterification using bromoalkyl mesogens as electrophiles. A similar approach was used by Pugh and Percec87 but this time the nucleophile was a carboxyalkylmesogen, the electrophile being a pendant chloromethyl group of poly(epichlorohydrin).


250·C

RONa

Cl

I +N=P+

I Cl

RONa

OR

I +N=P+

I OR

Fig. 7.25. Synthesis of phosphazene LCPs.

If PTC is not used, then a polar aprotic solvent is essential to solubilize the ionic nucleophile well enough for efficient reaction to occur. Chen88

has used hexamethylphosphoramide as solvent for the esterification of poly(sodium acrylate) with bromoalkyl mesogens while Pugh and Percec used dimethyl formamide (DMF) for the etherification of sodium 4-methoxybiphenoxide with poly(epichlorohydrin).21 In a compromise between PTC and single-solvent approaches, Maa and Chen have reported the efficient esterification of acrylic polymers with bromoalkylmesogens in DMF after isolating the tetrabutylammonium salts of the polyacrylates89 (tetrabutylammonium compounds are common PT catalysts).

7.5.5 Hydrosilylation of Alkenes In the hydrosilylation of alkenes, the Si-H bond, e.g. of a poly(hydrogen methyl siloxane), adds across the double bond of an alkene. The reaction, illustrated in Fig. 7.26, is normally catalysed by platinum(II), e.g.


Me Me I

+Si-O+ + 7'R I

H

I +~o+

Pt(ll> ~

R

Fig. 7.26. Hydrosilylation of an alkene.

391

chloroplatinic acid or dichloro(endodicyclopentadiene)Pt(II), and is of wide scope, but great care must be taken with the quality of the catalyst otherwise side reactions are probable, especially cross-linking through cyano groups in relevant mesogens (R). The reaction has been comprehensively reviewed by Gray!! who emphasizes the pitfalls associated with, for example, inattention to catalyst quality, incomplete reactions and poorly purified products. Using good catalyst the reaction can be effectively monitored spectroscopically, e.g. by the disappearance of Si-H stretching in IR spectra when excess vinyl mesogen is used. In such preparations complete removal of excess mesogen is essential; Gray and coworkers have demonstrated significant differences in both the phase-transition temperatures90 and the molecular dynamics of alignment phenomena9! for samples of the same siloxane LCP after different numbers of reprecipitations. Pure siloxane comb LC oligomers of relatively low molecular mass can be obtained using supercritical fluid purification.92

Gray!! describes several hydrosilylation reactions used to prepare side chain siloxane homo- and copolymers using well-characterized starting polymers, and Finkelmann 74 has modified the procedure for network siloxane LCPs by including bisvinyl-terminated oligosiloxane crosslinkers as well as vinyl-terminated mesogens in the hydrosilylation feedstock. But the method is also applicable to reactions that do not involve siloxane main chains. Thus Zen tel's 73 polymalonate elastomer was cross-linked using hydrosilylation of the terminal vinyl groups on the mesogens (Fig. 7.24) with oligo(dimethyl siloxane) cross-linkers bearing Si-H functionality at each end. Hydrosilylation has also been used to attach mesogens to the pendant vinyl groups of polybutadiene via spacers terminated by an Si-H group.93

7.6 APPLICATIONS AND APPLICABLE MATERIALS

7.6.1 Electro-optical Applications While comb LCPs exhibit display-relevant properties, they do not seem appropriate for real-time active displays of the type associated with


SMLC; the viscosity of the polymer matrix imposes switching times that are too slow.94 Instead, use can be made of the combination of thermoplastic and thermotropic behaviour. Nakamura et al. 95 have listed the merits of polymeric versus metallic components, and of polymeric versus SMLC components, in electro-optic devices. Compared to metals, polymers have low thermal conductivity, low toxicity, low manufacturing costs and good environmental stability. Compared to SMLC systems they offer possibilities of large size, free-standing, flat and thin displays, and may be compatible with directly deposited electrodes (avoiding cumbersome connection substrates). Even so, McArdle96 has pointed out that many practical problems remain before significant commercial exploitation becomes a reality.

7.6.1.1 Optical data storage The uses of comb LCPs as optical data storage media have been thoroughly reviewed.97 The modus operandi of a thermo-optic storage system is shown in Fig. 7.27.95 A homeotropically aligned transparent polymer at ambient temperature (a) is exposed to localized, disruptive heating (e.g. by a laser beam). The affected area is heated into the isotropic state with consequent loss of ordering (b), so that after exposure a laser track is recorded that is optically scattering (c). Storage may be short-term at ambient temperatures, although ambient smectic phases offer reasonable longevity.98 Durable storage can be achieved via vitrification after quenching of the exposed polymer to below Tg • Devices can

Laser beam

Fig. 7.27. Scheme for optical data storage. (After Ref. 95.)

be envisaged that offer permanent data storage (write-once devices), or erasable storage (write-erase devices) via reheating of the entire display and realignment (Fig. 7.27 (c)~(a)).

Various designs are possible to allow for choice of contrast (e.g. dark track on light background or vice versa), and devices can be fabricated from LCPs doped with dyes for coloured storage or for enhanced


brightness. However, doping with dyes compatible with the mesogenic units has another benefit; the efficiency of absorption of laser energy is greatly enhanced, allowing 'writing' with low pulse energies on a short timescale. Dye moieties can be present as guest molecules in guest~host systems (e.g. based on homopolymers blended with smallmolecule dyes), or may be chemically bonded onto the backbone in copolymer systems. The literature cited97 contains several examples based on both nematic and smectic polymers of types encountered earlier in this chapter.

7.6.1.2 Cholesteric LCP devices The thermochromic properties of a cholesteric mesophase were mentioned earlier (Section 7.2.3), and various devices are possible using cholesteric LCPs. Thermal-head recording of storable data onto a siloxane copolymer with cholesteric and nematic side chains has been describedY5 The copolymer was designed to display selective reflection of visible light and the write-on mechanism was similar to that shown in Fig. 7.27, but a thermal head was used to impress data onto local regions of the polymer (instead of laser addressing). Disrupted regions displayed good scattering contrast onto the reflective background, could be stored as required by rapid cooling and erased by reheating and slow cooling.

Brauchle and co-workers have reported two devices based on cyclosiloxane backbones containing cholesteryl and other substituents as shown in Fig. 7.28.99 The system n = 6~20, R = R 1, R 2, R 3 provides a device for high-contrast, write-once storage using a simple semiconductor laser, while materials based on n = 4 are appropriate for reversible, holographic storage with different selective reflection wavelengths depending on which azo-mesogen (R4 or R5) is present along with R 1 and R 2. The selective reflection associated with cholesteric phases can be vitrified, producing materials for use as selective filters or reflectors. 59

7.6. 1.3 Ferroelectric comb LCPs Ferroelectric behaviour in a material implies a macroscopIc dipole. In most materials that contain polar molecules, the dipole moments associated with the macroscopic sum of constituent molecules are non-additive; each molecule is dipolar but the molecules are so arranged in the material that the moments average to zero in the absence of an external polarizing influence. The key property of a ferroelectric

394

Me-Si-R I o

n

where

R = R', R2, etc.

Derek J. Simmonds

Fig. 7.28. Cholesteric cyclosiloxanes for optical data storage.

material is the possession of a spontaneous polarization that is nonvanishing in the absence of a polarizing field.

The symmetry elements of a tilted smectic C mesophase reduce to one symmetry axis (C2 ) when the smectogens are chiral (giving a chiral smectic C* mesophase). A single layer of such a smectic phase, possessing low symmetry, is capable of exhibiting a permanent dipole moment along the C2 axis.loo However, in chiral mesophases the director, as we have seen, undergoes a helical displacement from layer to layer. Each complete wind of the director helix would then average out the dipole to zero since a 3600 rotation of a mesogen's director will also correspond to a 3600

rotation of the C2 axis. But Clark and Lagerwall have described 101 a LC cell that is thin relative to the pitch of the director helix and that permits anchoring of the meso gens onto the cell surfaces. This surface-stabilized ferroelectric liquid crystal (SSFLC) device allows the helix of a chiral mesophase to be unwound, producing a sample with a macroscopic dipole.

A particular S~ phase, at a particular temperature, is associated with a particular tilt angle, e, to the direction of the polarization vector but, in the absence of an applied field, the two possibilities, + e and - e, are degenerate. However, exposure of the ferroelectric to an applied field


removes the degeneracy and switching to the lower energy tilt occurs if the field strength is above a threshold value. The tilt ordering remains after the field is switched off (a usable memory effect). Required field strengths are low, optical contrast between unswitched and switched states is good, and switching speeds are very fast-orders of magnitude faster than for twisted nematic cells1oo-and hence there is intense interest in this phase. Potential applications of LMMLC ferroelectrics have been reviewed by Goodby,lOZ while Le Barny and Dubois have discussed chiral smectic C comb LCPs. Z5

Some acrylate combs with chiral smectic C mesophases are shown in Table 7.14. Several related polymers (e.g. R=Rz, X=H, n=6, 11; R=Rz, X=Me, n=2, 6; or X=CI, R=Rz, n=2, 6, 11) did not exhibit S~

phases. 103 The biphenyl comb 104 was one of several similar products of different DP. The others exhibited similar meso phase profiles but with transition temperature variations consistent with the trends noted in

x

H H Me Me

R

Table 7.14 Chiral smectic C acrylate combs

X

I +CH2-C+

I *' COOCCH2)nO-R-OCH2CHCMelEt

o

R3 = --@-CH=CH-~-O--@-

n Phase transition temperatures (K)

109 284 S, 325 S~ 330 SA 389 i (cooling only) 2 g 338 S~ 383 SA 419 i

11 K331S~363SA379i 11 g308K338S~388SA415i


Section 7.3.1.2. Switching response times were shown to be shorter for polymers of lower DP but the variation levelled off at DP ~ 30. Each polymer switched most quickly at a temperature near the upper end of its S~ phase, and switching times were around 4-15 ms-much faster than those associated with nematic comb LCPs.

Ferroelectric combs, if realizable in device form, may have valuable applications; Le Barny and Dubois25 discuss their potential in the areas of display devices, transducers, pyroelectric detectors and non-linear optical devices (see below).

Ferroelectric siloxane combs have also been prepared (cf. Ref. 25), and some examples are shown in Fig. 7.29. Polymers of type I (n=6, 10, 11) all gave room-temperature S~ mesophases while polymers of type II (n = 6, 8) gave no S~ phase at any temperature, highlighting the importance of the constitution of the spacer chain.lo5 The length of the spacer affected both the detailed transition temperatures (cf. Table 7.14) and the switching response times, which decreased as n increased and which were well below 1 s at appropriate applied voltages.

o ° ~II 011 '" (I) R= -o-C-0---y-C-OCH2CH(MelEt

x=25,36,80

Fig. 7.29. Ferroelectric siloxane combs.


Polymers of type III 106 were crystalline and exhibited only SA phases when n =4 or 6 but gave S~ phases (from 363 to 378 K), independent of DP, when n = 10; long spacers seem necessary for S~ phase formation. The significance of DP (i.e. x=25, 36 or 80) was that the apparent viscosity of the S~ phase decreased as x decreased; this diminished viscosity would presumably lead to faster switching times as observed for the biphenyl combs shown in Table 7.14 and discussed above.

Some fascinating heterocyclic S~ combs were prepared by Hahn and Percec55 as a part of their investigation of conformationally mobile mesogen systems (Fig. 7.30). All three polymer types exhibited phases which on preliminary examination appeared to correspond to S~, but the detailed phase profiles varied (as would be expected). Also of interest is the biphasic behaviour of copolymers based on the dioxaborinane system, i.e. polymers III. While the homopolymer (III, x = 80, Y = 0) was glassy (Tg decreased in the order I> II > III, and I also gave a melting transition), the copolymer (III, x = 13, Y = 27) showed two Tg's and the copolymer (III, x = 5, Y = 25) gave two Tg's and two melting transitions. The authors55 propose that these biphasic data correspond to transitions derived from independent motion of the backbone and from cooperative but independent motion of the mesogenic units, and they go on to suggest that the rarely encountered biphasic/monophasic phenomenon requires further investigation.

Fig. 7.30. Chiral combs with heterocyclic mesogens.


7.6.1.4 Non-linear optical effects When light propagates through a material a polarization, P, is induced as described by a relationship of the form:

P=Po+ Xl. E+ Xl. E· E+ X3• E· E· E· ...

where Po is the spontaneous polarization of the material, E is the tensor of field strength of the electrical component of the optical field, and the X terms are the optical susceptibilities of the material (first-order, secondorder, third-order etc.).

In linear (first-order) optics a simplified expression, Pcx::xE, is an adequate description and the propagation of light through such a material is characterized for example, by the independence of refraction on light intensity and the absence of any change in frequency of the light.

Materials that exhibit non-linear optical (NLO) properties have considerable potential in optical communication and information processing technologies, and organic materials of this type are the focus of much attention.107.108 In particular, NLO comb polymers are of interest as frequency modulators, film lenses and materials for the reversible storage of optical information.l09 The polymers have advantages (e.g. coating possibilities, compatibility with integrated circuit components, sizable non-linear coefficients and fast NLO response) as well as drawbacks (e.g. thermal stability and purity requirements) in comparison with inorganic and small-molecule organic competitors. I 10

Polymer design for NLO potential can be guided by consideration of the expression

/Jind = C( • E + p. E· E + ')'" E . E· E· ."

where /Jind refers to the induced molecular dipole moment, C( is the linear polarizability, p is the hyperpolarizability, y is the second hyperpolarizability, and so on. The expression is analogous, but on a molecular level, to that given above for the macroscopic polarization, P, induced in a material. Thus a second-order NLO material (a X2 material), appropriate for frequency doubling applications, for example, implies involvement of molecular species having a dipole moment dependency upon a substantive hyperpolarizability, p.

Even-order non-linear coefficients vanish if individual molecules possessing substantive p coefficients pack into a material structure that has a centre of symmetry; the induced molecular dipoles cancel out in the material, which therefore has negligible X2 . The comb LCP architecture has the advantage (compared to small-molecule materials) that NLO


components can be tethered at one end to a spacer linkage and then to the backbone, thus reducing their freedom to order centrosymmetrically. Furthermore, molecular entities with large f3 values are to be found among, for example, stilbene derivatives that are highly compatible with tooth components of more familiar comb LCPs. Thus comb copolymers involving LC and NLO tooth components (offering poling of the material by cooperative interaction of the compatible teeth) are both synthetically feasible (cf. Sections 7.4 and 7.5) and structurally viable as NLO materials. Furthermore, homopolymer combs that include, as teeth, NLO components such as nitrostilbenes are quite similar to many comb LCP structures that we have shown in this chapter, and ferroelectric combs are also associated with NLO properties.

A variety of comb polymers and copolymers with second-order NLO potential have been prepared. 111,112 The structures in Fig. 7.31 demonstrate structural parallels between small-molecule X2 materials such as N,N'-dimethylaminonitrostilbene (DANS) (I) and the polymerizable nitrostilbene derivative (11).112 This and other monomer and polymer systems are under active study as applicable materials at present.

(Ill

Fig. 7.31. Materials for non-linear optics.

7.6.2 Liquid Crystal Elastomers 1 !3

Structural and synthetic aspects of these materials were discussed earlier (Section 7.4.3), and reviews are available from key workers in the field. 74,75 The crucial phenomenological bonus offered by the network architecture is due to stress-optical effects. Thus, partial or macroscopic

400 Derek j. Simmonds

ordering or disruption of the LC moieties can be brought about by stretching or compression. Ordering effects could lead to applications such as ordered elastomer films of useful dimensions that are orientable by mechanical rather than electric field methods. Disruptive effects, involving, for example, the impression of disordered tracks onto a pre-ordered elastomer by mechanical stressing using a stamp, could lead to the preparation of elastomer films containing light-conducting channels with possible applications in information technology.114

ZenteF5 has drawn attention to the possibility of creating piezoelectric elastomers by the combination of ferroelectric comb architecture (Section 7.6.1.3) within a cross-linked system. The ferroelectric polarization of such a material should be modified by the application of mechanical force to which the macroscopic ordering is vulnerable. That would offer the possibility of transforming a mechanical signal into an electrical response, and the elastomer would behave as a piezoelement.

7.6.3 Chromatographic Applications Silicone polymers have a long history as stationary phases in gas~liquid (partition) chromatography (GC). Thus poly(dimethylsiloxane) and methylphenylsilicones provide the useful SE and OV series of phases. More recently, small-molecule LC phases 115 and their blends with OV materials l16 have been evaluated. The occurrence of ordering (e.g. nematic or smectic) in a stationary phase is itself beneficial but, in addition, different separation criteria seem to apply with LC phases. Whereas relative volatility of analyte components dominates separations on ordinary SE or OV columns, molecular shape becomes significant when LC phases are employed (e.g. in the resolution of the isomeric benzo[a]pyrene and benzo[e]pyrene).117

A major drawback with some GC stationary phases results from their own volatility at operating temperatures, which, over a period of time, threatens column integrity via column bleed. Small-molecule LC packings are vulnerable to this problem but the comb LCP architecture, involving both LC moieties and siloxane polymers, is an attractive solution since it incorporates chemical tethering of the mesogens to the polymer.

Since the pioneering work of Laub and co-workers 116 a variety of mesomorphic polysiloxane GC solvent phases (the so called MEPSIL solvents) have been prepared and evaluated (reviewed in Refs 117~119). They are copolymer combs, often including dimethylsiloxane segments as well as mesogen-bearing units. Many of them still suffer from column


bleed at elevated temperatures, but developments involving materials with some unsubstituted Si-H sites (for possible attachment onto inert packings or column walls) have been advocated. 120 Cross-linking of the stationary phase is another strategy leading to decreased volatility, although analyte retention times tend to increase with the degree of cross-linking.121

In summary, useful GC stationary phases based upon comb LCPs are available. Areas of active development involve suppression of column bleed, depression of the glass-to-LC transition temperature (important for volatile analytes) and broadening of the mesophase temperature range. The latter point favours nematic systems, since that phase tends to have better longevity. In parallel with these developments, investigations are under way to apply LC stationary phases to capillary GC,119 to supercritical fluid chromatography,122 and to liquid chromatographyYs

7.6.4 Miscellaneous Recently attention has been paid to the incorporation of additional activity into comb LCPs using either copolymerization or combined (main chain and side chain) systems. For example, the polymer shown in Figure 7.32 is a copolymer involving meso genic cyanophenylbenzoate entities and photoactive spiropyran units (A) that are reversibly convertible to merocyanine units (B) under the influence of heat or light. 123 The polymer is photochromic and thermochromic, giving rise to yellow, blue, or red forms that are interconvertible depending on the thermal treatment or phototreatment applied. Photosensitive thermotropic LCPs of this type are of interest in the development of new photoimaging technology.

Fig. 7.32. Photo- and thermochromic LCPs.


Electrical activity in polymer materials is a keenly sought property with potential applications in electrical and electronic engineering and in microelectronics. Non-centrosymmetric ordering of polar mesogenic units might induce piezoelectric properties (cf. Sections 7.6.2 and 7.6.3). In most cases comb LCPs are aligned using a.c. fields but recently it was demonstrated that application of a pulsating current (generated by rectification of the a.c. field) to, for example, a cyanobiphenyl methacrylate comb derivative produced on oriented film that was piezoelectric. 124 Impact-generated voltages compared favourably with established organic and inorganic piezoelectrics.

Finally, a quite new series of comb polymers has been reported (Fig. 7.33)125 that are both mesomorphic (biphenyl aryl ester side chains) and, after doping, electrically conductive (conjugated main chain). The polymers are made by heating complexes formed between the corresponding nitrile (R-CN) and a transition-metal halide, and the process can be carried out at a mesophase temperature to allow ordered polymerization. The relationships between micro- and/or macroscopic ordering and electrical conductivity have yet to be established, but fascinating thermoelectrical applications can be envisaged.

+C=N+" I R

Fig. 7.33. Electroactive mesomorphic polymers.

ACKNOWLEDG EM ENT

The author is indebted to K.A. Foster, Division of Chemistry, for his help with the tables and figures.

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13. 77. Gleim, W. & Finkelmann, H., Makromol. Chern., 1987, 188, 1489. 78. Hammerschmidt, K. & Finkelmann, H., Makromol. Chern., 1989, 190, 1089. 79. Mitchell, G.R., Davis, F.J. & Ashman, A.S., Polymer, 1987,28,639. 80. Broer, D.J., Finkelmann, H. & Kondo, K., Makromol. Chern., 1988,189, 185. 81. Broer, D.J., Mol, G.N. & Challa, G., Makromol. Chern., 1989,190, 19. 82. Rodriguez-Parada, 1.M. & Percec, V., J. Polym. Sci. Polym. Chern. Ed., 1986,

24, 1363. 83. Percec, V., Makromol. Chern., Macromol. Symp., 1988, 13/14,397. 84. Reck, B. & Ringsdorf, H., Makromol. Chern. Rapid Commun., 1985,6,291.

Griffin, A.C., Bhatti, A.M. & Hung, R.S.L., in Ref. 51. 85. Keller, P., Macromolecules, 1984, 17, 2937. 86. Keller, P., Mol. Cryst. Liq. Cryst., 1988, 155, 37. 87. Pugh, C. & Percec, V., Macromolecules, 1986, 19,65. 88. Chen, S.H. & Maa, YF., Macromolecules, 1988,21, 2697. 89. Maa, YF. & Chen, S.H., Macromolecules, 1989,22, 2036. 90. Nestor, G., White, M.S., Gray, G.W., Lacey, D. & Toyne, K.J., Makromol.

Chern., 1987, 188,2759. 91. Attard, G.S., Moura-Ramos, J.J., Williams, G., Nestor, G., White, M.S.,

Gray, G.W., Lacey, D. & Toyne, K.J., Makromol. Chern., 1987, 188, 2769. 92. Krishnamurthy, S. & Chen, S.H., Makromol. Chern., 1989, 190, 1407. 93. Robert, P., Villenave, J-1., Fontanille, M., Gilli, J-M. & Sixou, P., Mol.

Cryst. Liq. Cryst., 1988, 155, 161. 94. Attard, G.S. & Williams, G., Nature, 1987,326, 544. 95. Nakamura, 1., Veno, 1. & Tani, c., Mol. Cryst., Liq. Cryst., 1989, 169, 167. 96. McArdle, c.8., in Ref. 1, p. 357. 97. E.G. Coles, H.J., J. Chern. Soc. Faraday Discuss., 1985,79,201. Coles, H.J., in

Developments in Crystalline Polymers, ed. Bassett, D.C., Elsevier Applied Science, 1988, Vol. 2, p. 297; cf. Refs. 62, 95 and 96.

98. Coles, H.J. & Simon, R., in Ref. 5b, p. 351.


99. Ortier, R., Bruchle, c., Miller, A. & Riepl, G., Makromol. Chem. Rapid Commun., 1989, 10, 189.

100. Walba, D.M., Slater, S.c., Thurmes, W.N., Clark, N.A., Handschy, M.A. & Supon, F., J. Am. Chem. Soc., 1986, 108, 5210.

101. Clark, N.A. & Langerwall, S.T., Appl. Phys. Lett., 1980, 36, 899. 102. Goodby, lW., Science, 1986, 231, 350. 103. Esselin, S., Noel, c., DeCobert, G. & Dubois, lC., Mol. Cryst. Liq. Cryst.,

1988, 155, 371. 104. Uchida, S., Morita, K., Miyoshi, K., Hashimoto, K. & Kawasaki, K., Mol.

Cryst. Liq. Cryst., 1988, 155, 93. 105. Suzuki, T., Okawa, T., Ohnuma, T. & Sakon, Y., Makromol. Chem. Rapid

Commun., 1988, 9, 755. 106. Keller, P., Ferroelectrics, 1988,85,425. 107. Chemla, D.S. & Zyss, l (Edsj, Nonlinear Optical Properties of Organic

Molecules and Crystals, Vol. 1. Academic Press, 1989. 108. Hahn, R.A. & Bloor, D. (Eds), Organic Materials jar Non-linear Optics.

Royal Society of Chemistry Special Publication No. 69, 1989. 109. Wendorff, J.H. & Eich, M., Mol. Cryst. Liq. Cryst., 1989,169,133; cf. Ulrich,

D.R. & Mohlmann, G.R. in Ref. 108, p. 241 and p. 275, respectively. 110. For a recent review see Mohlmann, G.R. & van der Vorst, C.P.J.M. in Ref.

1, p. 330. 111. E.g. Griffin, A.C., Bhatti, A.M. & Hung, R.S.L., Mol. Cryst. Liq. Cryst., 1988,

155, 129. Noel, C., Friedrich, c., Leonard, V., LeBarny, P., Ravaux, G. & Dubois, J.c., Makromol. Chem., Macromol. Symp., 1989,24,283.

112. Griffin, A.C. & Bhatti, A.M., in Ref. 108, p. 295. 113. See Finklemann, H., Gleim, W., Hammerschmidt, K. & Schatzle, l, Mak-

romol. Chem., Macromol. Symp., 1989,26,67, and cited literature. 114. Finklemann, H., Angew. Chem. Int. Ed. Engl., 1988, 27, 987. 115. Janini, G.M., Adv. Chromatogr., 1979,17,231. 116. Laub, R.J., Roberts, W.L. & Smith, C.A., J. High-Res. Chromatogr.,

Chromatogr. Commun., 1980,3,355. 117. Janini, G.M., Laub, R.J., Purnell, lH. & Tyagi, O.S. in Ref. 1, p. 395. 118. Witkiewicz, Z., J. Chromatogr., 1989, 466, 37. 119. Mazur, Z. & Witkiewicz, Z., LC-GC Int., 1990, 3, 38. 120. Janini, G.M., Laub, R.J., Pluyter, lG.L. & Shaw, T.J., Mol. Cryst. Liq.

Cryst., 1987, 153, 479. 121. Bradshaw, lS., Schregenberger, c., Chang, K.H-C., Markides, K.E. & Lee,

M.L., J. Chromatogr., 1986, 358, 95. 122. Rokushika, S., Naikwadi, K.P., Jadhav, A.L. & Hatano, R., Chromatog

raphia, 1986, 22, 209. 123. Cabrera, I., Krongauz, V. & Ringsdorf, H., Mol. Cryst. Liq. Cryst., 1988,155,

221. 124. Sato, K., Otsuka, K. & Matsumoto, M., Makromol. Chem., Rapid Commun.,

1988,9,631. 125. Barbarin, F., Dugay, M. & Fauxpoint, D., Mol. Cryst. Liq. Cryst., 1989, 167,

61.

Chapter 8

Thermotropic Main Chain Liquid Crystal Polymers

W.A. MacDonald ICI Wilton Materials Research Centre, P.O. Box No. 90, Wilton,

Middlesbrough, Cleveland, UK, TS68JE

8.1 INTRODUCTION

Low-molecular-weight liquid-crystalline compounds have been known for about 100 years.! However, main chain liquid crystal polymers (MCLCPs) have attained prominence only in the last 15 years. In 1956 Flory predicted lyotropic behaviour2 and this theoretical prediction was well demonstrated in the synthetic polymer area with the discovery by Kwolek of the aramids, e.g. poly( p-phenyleneterephthamide)? This led to interest in thermotropic main chain LCPs and, although aromatic polyesters which are thermotropic were described during this period in patents issued to ICI4 •5 and Carborundum Co} their liquid-crystalline nature was not reported. The first well-characterized description of a polymer exhibiting thermotropic behaviour appeared in the mid-1970s when Jackson 7 described a series of copolyesters prepared by the acidolysis of poly(ethylene terephthalate) with p-acetoxybenzoic acid which exhibited the phenomenon of opaque melts, low melt viscosities and anisotropic properties. These copolymers were test-marketed under the code X7G by Eastman Kodak but serious commercial interest did not develop, largely because the high-temperature performance of these materials was limited by the presence of the aliphatic component.

Wholly aromatic thermotropic polyesters were developed soon after by Celanese, who defined families of copolymers based on 2,6-naphthalene

407

408 WA. MacDonald

dicarboxylic acid, 2,6-dihydroxynaphthoic acid and 6-hydroxy-2-naphthoic acid,B.9 and by DuPont, who concentrated on copolymers containing ring-substituted monomers such as chloro-, methyl- or phenylsubstituted hydroquinone. 10. 11 These wholly aromatic polyesters had many of the same bulk and melt appearance characteristics as X7G, but had high-temperature performance characteristics.

Since the mid-1970s there has been considerable research in both industry and academia on liquid crystal polymers, with academic interest tending to focus largely on characterizing and understanding the structure~property relationship of these materials. The mid-1980s saw the launch of Celanese's Vectra and Dartco's Xydar (now marketed by Amoco) on the marketplace and there are now approximately 40 large companies known to be active in LCP research internationally, with further work being carried out in smaller specialized companies. 12 This chapter will describe the synthesis and properties of MCLCPs and will attempt to demonstrate the intimate interaction between the polymer architecture, morphology and bulk mechanical properties.

The following abbreviations are used in this chapter: HBA = 4-hydroxybenzoic acid; IA = isophthalic acid; TA = terephthalic acid; BP = 4,4'biphenol; HNA = 2-hydroxy-6-naphthoic acid; HQ = hydroquinone.

8.2 DESIGN OF THERMOTROPIC MAIN CHAIN LCPs

Completely rigid rod-like molecules such as poly(4-oxybenzoyl) or poly( p-phenylene terephthalate) tend to be highly crystalline and intractable, with melting points above the decomposition temperature of the polymers (>450°C). The problem of thermotropic MCLCP design is to disrupt the regularity of the intractable para-linked aromatic polymers to the point at which mesomorphic behaviour is manifested below the decomposition temperature and the materials can be processed in fluid yet ordered states. The disruption must not, however, be taken to the stage where conventional isotropic fluid behaviour is preferred. These requirements that the polymer must retain some rod-like nature but at the same time be melt-processable below 400-450°C have limited thermotropic MCLCPs mainly to polymers based on the linear ester or ester/amide bonds. With polyester/ polyesteramides, disruption is normally achieved by the three copolymerization techniques outlined in Fig. 8.1,13 i.e. frustrated chain packing, flexible spacers and non-linear links.

Typ

ica

l p

recu

rso

r h

om

o a

nd

co

po

lym

er

seq

ue

nce

s. In

tra

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ble

, T

m's

;;.

40

0°C

~@C1 (

O<B@

=) {o

(@).,

ot, lc

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~

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~Phati

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nk

s

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ate

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~:.~,c

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} (o

~o){

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)

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. 8.

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ermotro

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13

..., ;::; '" ... ;; c ~ '" r; .

~ '" s· r'l

;::;- '" s· t-- .E. '" is.: Q

~ '" §'.

'1:l

c ~

;; '" ~ -1'0 ~

410 W.A. MacDonald

8.2.1 Frustrated Chain Packing Frustrated chain packing refers to any mechanism which while maintaining essential linearity and chain stiffness makes close and regular correlation into a 3D lattice difficult. Copolymerization with linear unsubstituted phenyl-based residues to give a random copolymer does result in a reduction in Tm , but the reduction is not enough to render the polymers readily melt-processable. Instead three classes of disrupters are used. The first is the incorporation of a parallel-offset or crankshaft monomer such as a naphthalene unit. Units of this type have been found to be very effective at reducing Tm.14.15 Figure 8.2 shows plots of Tm vs molar percentage composition for two different copolyesters with and without crankshaft moieties which illustrate this point.

Fig. 8.2. Variation of Tm with composition for two co polyesters with and without a crank

shaft-like moiety.14.15

500

450

400

350

300 " HBNHNA /

250 ~ 200

:or iii i

20 40 60 80

Mole % p-Hydroxybenzoic Acid (HBA)

The second type of disrupter is based on a biphenyl unit which extends the monomer unit. This type of unit, which lacks a parallel offset, does not appear to be as effective as the naphthalene unit for depressing Tm. 6

The third type of disrupter is based on ring-substituted phenyl groups where chain packing can be frustrated by sequence randomization resulting from the random occurrence of head-to-head and head-to-tail isomerization if the unit is unsymmetricaPO.ll or by steric effects such as increased chain separation and decreased coplanarity of adjacent units in the meso gen. 16,1 7,18

Comparison of substituents of similar size but different polarity (CH 3 ,

Br, N02 and CN) with substituents of different size (H, CH3, CI and Br) demonstrated that the size of the substituent rather than the polarity was the important factor, with transition temperatures (Tg, Tm and isotropization temperature) decreasing with increasing substituent size in a homologous series of monosubstituted polymers. 17.18 This indicated that the predominant steric effect of the substituent was to increase the separation of the mesogenic units and reduce the molecular packing

Thermotropic Main Chain Liquid Crystal Polymers 411

efficiency. In general, however, ring substitution alone does not sufficiently lower the Tm to permit melt-processability unless large groups such as the phenyl are used. 19 Also, the ring substituent can adversely affect the polymer thermal stability, in particular with the CH3 and Cl substituents. 19

8.2.2 Polymers with Flexible Spacers Thermotropic behaviour is observed in many polymers that contain rigid-rod units alternating with flexible spacer units of the type (CH 2) and (CH2CH 20).20 The length and length: diameter ratio of the rigid-rod unit and the length of the flexible spacer unit affect the type of mesophase formed and the transition temperatures, with the nematic phase stability being highest with the greater the length and length:diameter ratio of the rigid-rod unit and the shorter the flexible spacer unit. The shortest rigid-rod unit so far reported to permit nematic melt formation is21

/o~d with a rigid-rod length of only about 1.1 nm.

An odd-even effect on the crystal nematic and nematic isotropic transition temperatures has been observed with increasing numbers of methylene units in the spacer group.22-26 In general, MCLCPs containing an even number of methylene units exhibit higher transition temperatures and entropy changes than those having odd units. A typical phase diagram illustrating this is shown in Fig. 8.3.24 This odd-even effect appears to be strongly related to the trans-gauche conformation of the flexible spacer influencing the stability of the mesophase. 27 Increasing the flexibility of the spacer group, for example by replacing a methylene with a siloxane spacer, significantly reduces transition temperatures. 2B

In addition to having rigid-rod units and flexible spacers regularly disposed, the situation can also arise where the rigid-rod length and/ or spacer length are no longer constant throughout the chain, for example in the MCLCPs based on copolymers of poly(ethylene terephthalate) and poly(4-oxybenzoyl), i.e. the X7G family of polymers.7 In polymers of this type there is a distribution of rigid-rod lengths; the effect of this distribution on the nematic --+ isotropic transition has not been established but the crystal--+ nematic transition of these polymers tends to be lower relative to the corresponding polymers with regularly disposed spacers, which tends to give a nematic phase

412

T/k

W.A. MacDonald

580

560

\ ok - Ic 540 \

• Ic- i I I

\ 520 I \ \ I

500 I I \ I

480 I I I I I \ I I I I

460 I I

'rf 440

420+---.---.----.-.--.-----.----, 7 8 9 10 11 12 13 14

n

Fig. 8.3. Variation of TKN and TN! with the number of methylene units for homopolyesters of 4,4'-dihydroxy-y-methylstilbene and aliphatic acids.24

of wider temperature range. 29 Over 30% of the oxybenzoyl units have to be incorporated before thermotropic behaviour is observed. 7

A range of copolyester systems incorporating ethylene terephthalate units and rigid-rod mesogenic units have been prepared 30 and the analogous polyesteramide systems can be prepared from poly(ethylene terephthalate) and p-acetamidobenzoic acid, although the presence of the amide links gives rise to a higher-melting polymer relative to the polyester equivalent. 31

8.2.3 Non-linear Links The third common strategy for lowering Tm in MCLCPs is to introduce non-linear units into the polymer backbone. These can take the form of rneta- or ortho-substituted phenyls, 1,6- or 2,S-linked naphthylenes 13 ,14,20,32 or the incorporation of kinked bonds within the monomer unit such as -C(CH3h-, -CH2-, -0-, -S-, -S02-, -CO-,20 This method of disruption is especially effective at lowering Tm, but the introduction of large amounts of such bent units will lead ultimately to loss of liquid crystallinity. In general, the stereochemistry of the bisphenol containing a kinked bond is more important than the polar


effects imparted by the type of substituent connecting the two phenolic rings in the bisphenol. Lenz et al. showed that large substituents such as -C(CH3)z- which cause more deviation from linearity and copolarity of aromatic rings more effectively disrupt the rigid rod-like nature of the polymer. The effectiveness of units of these types in disrupting the polymer is ranked as shown in Table 8.1. 20

Table 8.1 Effect of amount of bisphenols on the properties of MCLCPs based on chloro

hydroquinone and terephthalic acid

Structure of bisphenol

CH3

HO-@-9-@--OH CH 3

HO-@-~-@-OH o

HO-@-CH2-@-OH HO-@-S-@-OH H0l§r0H

HO-@-O-@-OH HO-@-@-OH

Maximum mol % in copolyester for retention of liquid crystallinity

40

50

60

60

60

70

100

f~j--@-~l f--@-x--@-oJ--@-g} x= none, -C(CH3J2-, -CH2-, -0-,-5-, and -502-

Kinks can also be incorporated into the backbone by preparing random copolymers which in addition to ester functionally also contain non-linear bonds such as anhydride 33 or carbonate bonds. 34,35 However, given that the ultimate mechanical properties will depend on how

414 W.A. MacDonald

rod-like the polymer is, the introduction of kinks will ultimately effect mechanical properties. Furthermore, the presence of meta-aromatic isomers such as isophthalic acid can have a detrimental effect on chemical and hydrolytic stability.14 The effect of structure on these properties will be discussed later.

Figure 8.4 illustrates the typical monomer unit types that are used to lower Tm 14 and Table 8.2 shows a range of typical MeLeps developed by industry which illustrate the various strategies for lowering Tm. This table is by no means exhaustive, but most other polymer formulations tend to be variations on the strategies shown. As can be seen from Table 8.2, combinations of the various strategies can also be used to lower Tm. From an industrial point of view the frustrated chain packing approach to achieving melt-processable MeLeps is preferred as this tends to give rise to polymers of the highest degree of liquid--crystallinity in terms of the inherent rodlike nature of the polymer chain, which will ultimately effect flow and

-OCH.CH.O-

Aliphatic [email protected]@-

[email protected]@-

Bent rigid ~

Swivel -@-x-@- ~ x= O.S.C

Parallel offset ~O 0 rcY0Yc)y "Crank Shaft" ~ ~!8J

° Ring substituted ;§ex

x= CI, CH3 .Phenyl

Fig. 8.4. Strategies for lowering Tm in MCLCPS.14

Table 8.2 MCLCPs developed by industry, illustrating the various strategies for controlling

Tm

Celanese

Dupont

Dupont

Owens Corning

( Montedison/) Granmont

Carborundum (Amoco)

ICI

o o II

(o~g)(o~C) o

(o~g)~©g)(o©o) o

(o~g)~©g)(o©N0

o 0

(o@-©o)(g@g)(o©J)(gcgi)

R - aromatic diol, carboxylic acid, hydroxy acid (less than 2t%)

Ref 8

9

36

10, 37, 38

11

39

40,41

6

42

416 W.A. MacDonald

the ability of the polymer to orient. These types of MeLeps also tend to have higher temperature performance than polymer containing kinks or flexible groups and better chemical and hydrolysis resistance.

8.3 SYNTHESiS OF THERMOTROPIC MAIN CHAIN LCPs

Some of the early MeLeps were made either by interfacial polymerization or by high-temperature solution polymerization from diphenols and dicarboxylic acid chlorides. This route gives rise to polymer where the rigid-rod units alternate regularly with the disrupter group and is the method generally used with polymers containing flexible spacers. The majority of MeLeps, especially those of commercial significance, are prepared by an ester exchange reaction between acetoxyaryl groups and the carboxylic acid group with the elimination of acetic acid, at temperature above the crystalline melting point of the polymer produced. 14 The bulk of the reaction and acetic acid evolution is accomplished at normal atmospheric pressures, after which the resulting oligomers are polymerized to high molecular weights by application of a vacuum as outlined in Fig. 8.5. Under

o II~

CH3CO\Q; C02H ~ ~C02H

CH3CO~

Inert I 200'C Gas ,

Clear Melt

%-3 Hrs ~ 25~ Acetic Acid Evolved

Turbid Fluid Melt

10 Min-1 Hr ~ 280-340°C Vacuum

Opalescent Polymer Melt

, Extrude

t~~,~WY, Fig. 8.5. Typical MCLCP melt polymerization synthetic scheme.


such conditions the polymer often forms a liquid crystal (LC) phase during polymerization and the melt is characterized by exhibiting shear opalescence. Condensation polymerization starting with the phenyl esters of the diacids with the aryl dials has also been reported. 6 Poly( p-oxybenzoyl-co-ethylene terephthalate) (X7G) falls into a special category within this general synthetic route. In this particular case the polymerization proceeds by cleavage of the molten poly(ethylene terephthalate) followed by condensation under vacuum of the carboxy-terminated and acetate-terminated segments, and also self-condensation of the p-acetoxybenzoic acid. 7 Inhomogeneity can be a problem with this particular synthesis.

High-melting polymers have been prepared by a condensation reaction in an inert heat-transfer medium6 .43 where the temperature of the stirred reaction mixture was raised very slowly over many hours. More recently a novel high-temperature non-aqueous dispersion polymerization route to MCLCPs has been reported. 44 Polymerization is carried out in an inert heat-transfer medium but also in the presence of polymeric stabilizers and/or hydrophobic inorganic stabilizers which sit at the surface of the polymer droplet and prevent flocculation during the comparatively rapid raising of the temperature over 1-2 hours to the final polymerization temperature. Low-melting MCLCPs (Tm < 340°C), which would normally agglomerate in the heat-transfer medium under such conditions in the absence of such stabilizer, can be prepared by this route and the molecular weight of the final polymer is not viscosity limited, as with the melt polycondensation reaction, by the requirement that the melt has to be extruded from the autoclave. The polymerization follows the route oulined in Fig. 8.6. Unlike the melt process the polymerization can be carried out below the melting temperature of the polymer and no vacuum is required to obtain high molecular weight because of efficient removal of the acetic acid byproduct from the small polymer droplets. This process also yields polymer in a unique unsheared form.

To get around the problem of the viscosity of the melt polycondensation limiting molecular weight, the polymer can be polymerized in the solid state by raising the temperature in stages to below the melting temperature. 7 This route can also potentially be used for obtaining high-melting polymers by preparing a low-molecular-weight prepolymer by the melt route followed by subsequent solid-phase polymerization to complete the polymerization.45

418 W.A. MacDonald

o II ~

CH3 C 0 Q; C02H +

~OOC02H CH3C~

Liquid Paraffin

Stabilisers Added - ,

Onset of Polymerisation Stirrer Speed Increase , Dispersion of Droplets

Droplet Coalescence to Equilibrium size Distribution Dependent on

Availability of Stabiliser

300"C , "Acetic Acid Evolved

High Molecular Weight Polymer within Stable Polymer Droplets

Cool, Centrifuge , Particles 1 0 - 150 J.Lm ,

Solvent Wash , Free Flowing Powder

Fig. 8.6. Typical MCLCP non-aqueous dispersion polymerization synthetic scheme.

Direct reaction of diphenols with dicarboxylic acids46 and 4-hydroxybenzoic acid with 2,6-hydroxynaphthoic acid47 in the presence of tin, titanium or antimony compounds has been reported. Although high molecular weights are reported, the reaction works less efficiently than when the phenolic groups are acetylated. The reaction involving 4-hydroxybenzoic acid has the added complication of decarboxylation of the 4-hydroxybenzoic acid competing with polymerization. Direct esterification is not used as a commercial process for the preparation of fully aromatic MCLCPs. MCLCPs other than polyester or polyesteramides have been reported. Recently several papers have described MCLCPs which are based on non-linear bonds which rely on a meso genic unit such as a biphenyl to impart enough rod-like nature on the polymer to give liquid-crystallinity. Percec


and Shaffer have prepared thermotropic polythioethers and copolythioethers48 based on 4,4'-dithiolbiphenyl and aromatic polyethersulphones49 by a two-phase (aqueous NaOH-KOH~organic solvent) phase transfer catalysed reaction. With the aromatic polyether sulphones several potential mesogenic units such as 2,6-dihydroxyanthraquinone, 2,6-dihydroxynaphthalene and 4,4'-biphenol were incorporated into the backbone, but only the 2,6-dihydroxyanthraquinone-containing polymer exhibited liquid-crystalline behaviour. Percec and Shaffer have also described thermotropic polyketones50 synthesized by the Friedel-Crafts arylation of biphenyl and fluorene with IX,w-dicarboxylic acid alkanes in the presence of a phosphorous pentoxide/methanesulphonic acid condensing agent.

The synthesis of thermotropic polyurethanes,5! polyethers52 and aromatic polyazomethines53 has been reported by other researchers, but at present research on MCLCPs of the type described in this section is concentrated in academia and there has as yet been no major industrial exploitation. Block copolymers of polyarylsulphones and ketones 54-56 are currently exciting interest in a number of laboratories. These have been prepared by synthesizing polyaryl sui phones or ketones with phenolic functionality, acetylating the end groups, and treating these functionali zed blocks like a diphenol in a conventional LCP polymerization process.

8.4 CHARACTERIZATION AND MORPHOLOGY OF THERMOTROPIC MAIN CHAIN LCPs

8.4.1 Solution and Melt Characterization MCLCPs are insoluble in common laboratory solvents and characterization in solution requires the use of relatively exotic solvents such as pentafluorophenol, trifluoromethanesulphonic acid and chlorophenol (often at high temperatures) or mixtures of solvents such as trifluoroacetic acid and dichloromethane. This has restricted analyses of the solution properties, especially with the fully aromatic MCLCPs. Calundann et al. have studied the intrinsic viscosity-molecular weight relationships of a wholly aromatic copolyester using GPC-low-angle light-scattering techniques.!4 They calculated the Mark-Houwink coefficient to be 0·98 for molecular weights between 5000 and 50000, suggesting a semi-rigid conformation in solution. They also reported that NMR studies indicated a random monomer sequence, and this

420 WA. MacDonald

has been observed in other MCLCP systems. 15 Light-scattering studies on poly(phenylhydroquinone-co-terephthalic acid)57 indicated that the persistence length of 10 nm was smaller than that expected for a thermotropic polymer having a nematic-isotropic transition at 450°C and smaller than the persistence length calculated for completely rod-like intractable polymers such as poly(hydroquinone-co-terephthalic acid). It was argued from examination of space-filling models that the carbonyl group adjacent to the pendent phenyl ring is displaced from the plane of the benzene ring. This destroys the partial double bond in that carbonyl group, leading to a more flexible chain.

Although extensive light-scattering studies on a range of MCLCPs have not been carried out, it seems likely that melt-processable MCLCPs will exhibit smaller persistence lengths relative to the intractable para-linked aromatic polymers, given the requirement that the perfect regularity of these intractable polymers must be disrupted to give thermotropic MCLCPs as described in Section 8.2.

MCLCP melts are characterized by being opaque and exhibiting the characteristic shear whitening effect associated with anisotropic melts.14 Under crossed polars the melts exhibit birefringence and the typical threaded Schlieren textures associated with domain texture in areas of high local orientation. This domain texture is equivalent to the textures seen with small-molecule nematic liquid crystals. 58.59

Combining hot stage microscopy and DSC analysis allows fairly complete phase diagrams to be constructed for polymers which exhibit an isotropic and liquid-crystalline phase, as shown in Fig. 8.7. 13 The characterization of meso phases is discussed in detail in Chapter 2.

8.4.2 Morphology and Structure of MCLCPs The hierarchy of structure in MCLCPs can be divided into three main categories as illustrated in Fig. 8.8, i.e. macro, micro and molecular morphology. In this description the macromorphology is taken to include the structure of fabricated articles (millimetre to 10-J,lm scale) and the micromorphology to include domain and banded textures (10 J,lm-O·5 J,lm). Molecular morphology is taken to represent the way MCLCPs chains pack and crystallize and the conformational changes that the polymer backbone undergoes as the temperature is raised (submicrometre scale). There is obviously overlap at the 'boundaries' of the three morphologies.


T Biphasic /' N-I

, / / / , /

ReglOn~ /' /

400 / / // DecomposItion

: IsotropIc Temperature :Melt I I I I

---- ---r--ff------- ~ / -----------;------e: 300 i ThermotropIc

~ .2 ~ .. : I Melt I I / I ' ' ~

Q) Co

E Q)

: / / If / 1//

f- 200

TA 0 15

IA 100 85

50

50 Composition

100

o

421

Fig. 8.7. Transition temperatures versus composition as the ratio of T A: IA is varied.' 3

8.4.2.1 Macromorphology Injection mouldings, films and fibres of MCLCPs exhibit complex morphology which depends on the flow/processing conditions. The macroscopic morphology of injection-moulded bars consists of three macrolayers: two skins with a core between.60•6l The skin macrolayer has a distinct structural gradient comprising three subdivisions from the surface inwards~a highly oriented top layer, fibrillar in nature; several oriented sublayers consisting of stacks of interconnected microlayers; and a less-oriented inner zone (Fig. 8.9). On injection moulding, MCLCPs fill the mould by the advancing front method (fountain flow). The elongationaljextensional flow in the advancing front produces high levels of molecular orientation in the direction of flow at the surface skin. The core, on the other hand, is subjected only to shear flow, which produces significantly less molecular orientation which may be perpendicular to the flow direction representing localised flow patterns. The skin to core ratio is

422 W.A. MacDonald

Molecular Morphoplogy Crystall inity IS·l0nm) Chain conformat ion IO.Snm)

Balow Tm Aboye Tm

/ ~ I Crislalilies

(aCUlar "'ulions transitions, Tg etc

Stillness/ temperatu re characteristICS Melti ng behaviour

Micromorphology -1 ·10.um· Opt ical appearance

~ ~~ '--- W-~,,-- -...;e. Domains Disclinations

Mouldings. extrusions(> mm)

8 L8yerl00.u m

Skin 40% " Core 20%/ ";:'-==-_ -: ~:./ ~ Skin40% \ol ... _

Direction of Ilow

Macromorphology

Anisotropy, mechanical properties

I<' ig. KIt Hierarchy of structure in \1CLCPs.

approximately 2: 1 but depends upon the MCLCP's flow characteristics and processing conditions used.

8.4.2.2 Microstructure Extensive characterization of MCLCPs has resulted in the delineation of a fibrillar, hierarchical structure model which describes the microstructures that Jaffe and Sanger observed in a broad range of oriented fibers, extrudates and moulded articles. 62 Three distinct fibrillar species are observed: microfibrils that are about 50 nm, fibrils about 500 nm and macrofibrils that are about 5 /lm in size.

Polarized-light microscopy and transmission electron microscopy studies62 66 of MCLCPs which have been subjected to deformation such as shear have revealed the presence of a banded texture which appears as regularly alternating light and dark bands whose long axes lie perpendicular to the shear direction. The period is usually of the order 1-10 /lm, but this is known to depend on such variables as shear rate and temperature. Unlike the lyotropic polyphenylene terephthalamide, where the molecular organization underlying the banded texture appears to correspond to sharp changes in the molecular orientation,67.68 with thermotropic MCLCPs the

Skin Core Skin

Sublayers


Top layer

Less ordered microlayers

1 Fibrous con microlayers

nactions between (diameter = O.2.urn)

o T 20~m

/SUblayer I IjiCrOlayer 'l Width ranges from ten to a few hundred micrometers

I

j/ t 1 " ... ~ 30 to 5O.urn

500 to 700 ~m

~ajor flow direct ;::>

ion

-?,um

T ~ay'!'_ f

423

Fig. 8.9. Baer et a/.'s schematic illustration of proposed hierarchical model in MCLCP injection-moulded bars.

changes in molecular orientation appear to be more gradual and smooth.

Optical studies in polarized light of a uniaxially oriented MCLCP fiber or thin tape revealed textures that have been described as domain-like, with the domains observed in these oriented structures being about 0·5 Jim across and elongated along the fiber axis. 62 Again, one possible explanation proposed for the colour variation between domains is that it is due to the serpentine trajectory of the molecules within a domain, compared to adjacent domains.

As yet there seems to be no suitable model to describe the formation of the bands, although a consensus does seem to have been reached that they result from some kind of relaxation phenomenon after the deformation has ceased. Factors that are likely to playa role in determining the two key parameters in describing the bands- their period and maximum deviation angle-are chain length, stiffness and the different Miescowiz viscosity coefficient, as they will determine what mode of relaxation is likely to dominate. However, there are very little data available on these relevant quantities.

From the periodicity of the bands it is clear that they correspond to some kind of supermolecular structure, with many molecules changing orientation systematically and in concert. Such macroscopic properties of MCLCP as the modulus will depend on the overall level of

424 W.A. MacDonald

orientation achievable. The presence of bands indicates that the orientation is regularly changing, and this is likely to have a significant effect on modulus. The area of bands and domains and the interplay with bulk mechanical properties and molecular morphology is an area that is currently poorly understood but is clearly important. One may hope that future research will continue to address this problem.

8.4.2.3 Molecular morphology The molecular morphology is the underlying key to the behaviour of MCLCPs and dictates to a large extent the macro- and micromorphological behaviour of MCLCPs.

The nature of crystallites in solidified rigid-chain MCLCPs of the types shown below (II and III) was examined by Blundell. X-ray diffraction and

f°-«tJ-@--11L-©-oJ-©--o-c",_c,,-o-©--11II J~5l ~5

ill

thermal analysis showed69 that identical heats of fusion were observed for slow-cooled specimens with well-developed crystals and for quenched specimens with almost amorphous X-ray patterns. This could only be explained by the presence of microcrystals that were much smaller than those possible in conventional polymers. If the surface energy per unit area were of the same value as that usually found in chain-folded lamellar crystals, then from

A Ilh=llh --")'

00 V (8.1)

where t:.hoo = heat of fusion per unit volume of an infinite crystal, A = surface area surrounding regions of three-dimensional crystals, )' = surface energy per unit area, and V = volume of crystals, the term (A/V))' would be so large that the crystal would be unstable and would spontaneously melt. It was proposed by Blundell that the surface energy of MCLCPs was small and a direct consequence of the molecular morphology of the nematic liquid-crystal state. Unlike with the vast increase in disorder experienced when a conventional lamellar crystal


Below Tm Above Tm

a

b

Fig. 8.10. Blundell's schematic diagram of the morphologies above and below the crystal melting point for (a) rigid-chain nematic polymer, and (b) conventional polymer with chain-folded lamellar crystals. The thicker parts of the lines

represent regions where the chains form 3D crystal lattices.69

melts, as illustrated in Fig. 8.10, it was envisaged that in a nematic liquid crystal that there is little change in the general configuration of the molecules before and after melting. This has important implications for warpage and shrinkage on moulding and will be discussed later.

It was also observed in the same paper that the heat of fusion (ll.Hr) and entropy of fusion (ll.Sf) of the MCLCPs studied were significantly lower than those of a corresponding polymer such as poly(ethylene terephthalate). This has important implications in the design of MCLCPs. The low ll.H f is directly related to the lower level of molecular cohesion within the crystallites, which results from chain irregularities, whereas the low ll.Sf is a direct consequence of the chain stiffness. Chain irregularities and chain stiffness are two distinct properties that can be separately designed into the MCLCPs. In the polymers of type III it automatically follows that there will be a low entropy change, since there is little or no change in overall configuration on crystallization. From eqn (8.2),

(8.2)

426 W.A. MacDonald

given that i1.Sf is low then the i1.Hf must also be low, otherwise Tm would be too high and the system intractable. Thus, in order to make processable MCLCPs based on frustrated chain packing, irregularities must be introduced into the chain in order to limit the effective bonding of the crystals. This then gives rise to the disordered crystals of poor 3D order.

There has been considerable research directed at elucidating the nature of the crystallites in MCLCPs of the above type. X-ray diffraction has shown that the polymer consists of random sequences of the monomer units only constrained by the necessity for chemical bonding. 7o,71 However, the presence of sharp X-ray diffraction maxima as well as the aperiodic meridional intensity distribution suggests that there is a certain amount of 3D ordering despite the random monomer sequence. This has been considered in two similar but different approaches by two groups. Blackwell et al.72 have considered that many adjacent molecules have a single plane of register in which identical monomers are side by side and that the register vanishes as one moves away from that plane. Recognizing the physical improbability of this approach, the rigid register plane was replaced by an out-of-plane distribution function for the monomers about the plane. It was found that this reproduced the diffraction pattern quite successfully with a distribution not differing much from random. An alternative model has been introduced by Windle et al. 73 - 75 which differs from that of Blackwell in one key respect. The interchain correlations are seen as being between adjacent similar sequences of monomer units which in general will be aperiodic. These associations occur by chance in a rapidly cooled sample, and statistical modelling shows that they are able to account for the observed crystallinity of ~ 10%. The matched sequences form small plate-like ordered entities and have been called non-periodic layer crystallites. Their predicted distribution and dimensions are consistent with diffraction and dark-field TEM observations.75 •76 The model is also extended to account for the improvement in crystalline order on annealing or slow cooling. It envisages interchain diffusion leading to segregation and thus more extensive matching of the sequences. Such predictions are confirmed by diffraction analysis, SEM observations of etched specimens77 and TEM (Fig. 8.11).

If the situation where the MCLCP is based on a kinked structure such as polymer IV, crystals with good 3D order are observed. The deduced values of i1.Sf and i1.Hf are considerably higher and are


Fig. 8.11. Scanning electron micrograph of NPL crystallites in an etched section of a random copolymer of type III (DP - 25). The crystallites are platelets normal to the molecular axis and in this view reveal a bend (tilt boundary

between two domains. (Photograph supplied by A. Windle.)

more akin to the values observed in conventional isotropic polymers such as PET 78 (Table 8.3). This can be rationalized in terms of the

meta links of the isophthalic acid giving rise to potential flexibility in the chain and hence enabling greater conformational freedom. If IlSr is larger, then from eqn (8.2) an acceptable Tm can be achieved with crystals of higher order and better chain packing, i.e. llHr can be larger.

Further evidence for the higher order in MCLCPs containing kinks comes from the effect of annealing on the density of MCLCPs based on frustrated chain packing versus MCLCPs based on kinked structures. It has been shown that LCPs based on polymer III show no change in density on annealing despite the crystallinity increasing. 79 Unpublished work from our laboratories agrees with the above results, but shows

428 W.A. MacDonald

Table 8.3 Crystallinity data on MCLCPs containing kinks or crankshaft units

compared to polyethylene terephthalate

Polymer X-ray ,1.Hr ,1.Sr fractional (kJ kg-I) (kJ kg-I CC I)

crystallinity

Polymer IV, 0·16 117 0·2 x=0·36

Polymer III, 0·25 20 0·04 x=0·73

PET 0·5 135 0·26

that an MCLCP based on polymer IV which contains isophthalic acid does increase its density as the crystallinity increases (as measured by wide-angle X-ray diffraction) (Table 8.4), i.e. the MCLCP containing kinks is achieving close regular packing in the crystalline regions and is behaving like a conventional isotropic polymer.

The implication of these observations is that the polymer architecture has a significant effect on the crystalline nature and this in turn has major implications for bulk polymer properties such as warpage and shrinkage (see later).

Table 8.4 Density versus crystallinity data on MCLCPs containing kinks or crankshaft

units

Polymer IV, x=0·36 Control Quenched Annealed 250cC for 24 h

Polymer III, x=0·73 Control Annealed 250°C for 24 h

Density (g cm- 3)

1·390 1·391 1·422

1·401 1·401

Percentage crystallinity

<5 <5 12·6

15 25

The ordered regions in MCLCPs appear to playa similar role to crystals in conventional polymers, i.e. they tie the molecules together and impart molecular rigidity. However, the level of crystallinity in MCLCPs is low, often under 20% for an unannealed sample, and the crystallites in MCLCPs account for only a minor proportion of the total material. The


remammg material, although conforming to a general nematic configuration, possesses no three-dimensional correlation between segments of adjacent molecules. The variation of stiffness with temperature in the region below the melting point of the crystallite depends very much on the molecular segment in the non-crystalline region. Taking LCPs based on frustrated packing, such as polymer III, the dynamical mechanical properties are typified by an appreciable fall-off in modulus with temperature which is particularly significant over the temperature range 0-120°C.80-82 Three separate molecular processes have been identified in M CLCPs of this type (Fig. 8.12). The upper 'X process at about 110°C appears to be related to the Tg transition of a conventional polymer, although the Tg is surprisingly low for a fully aromatic polyester. The prominent f3 process in the region of 50°C is associated with the motion of the naphthyl moiety, and the shallow y process at about - 40°C is attributed to similar motions associated with the phenyl unit. 14

MCLCPs based on kinks rather than frustrated chain packing do not exhibit the f3 loss process associated with the naphthyl moieties,

y

0.15 j

0.1

Tan 0

0.05

·100 ·50 o 50 100 150

Temperature 'C

{aU COt. t aijco t, (A) x =0.4 , Y = 0.6; (8) x = 0.7, Y = 0.3; (C) x = 0.8. Y = 0.2

Fig. 8.12. Tan t5 loss from dynamic mechanical analyses of copolymers of the form shown. 80

430 W.A. MacDonald

with the result that MCLCPs based on kinks, e.g. polymer IV (X = 0·36), have lost 50% of the O°C DMA modulus at 120°C (Fig. 8.13), whereas the naphthalene-containing MCLCPs have lost 50% of the O°C DMA modulus at 80°C.S1 However, although the MCLCP shown in Fig. 8.13 containing the kink has a superior modulus retention with temperature to the naphthalene containing MCLCPs, it can be seen that relative to the isotropic aromatic polyester analogue the MCLCPs still have a poor modulus retention with temperature. This rapid falloff in modulus in the 0-120°C temperature range is obviously undesirable in an engineering resin, but appears to be common to most MCLCPs. The reason for this behaviour is not clearly understood but may be related to large-range conformational motions occurring in the temperature regime 0-120°C and arising from the rod-like nature of the MCLCPs.

DMA of HBA,.IA"HQ3' Vs Polyarylate

Normalised DMA

Stiffness

0.5

o

I

./ '-'/ '-. . . . . . . . .

100 200 Temp °C

Tano

0.1

Fig. 8.13. DMA of a MCLCP containing kinks compared to an isotropic aromatic polyester. 81 (--) normalised DMA stiffness HBA36 IA32 HQ·32; (_._) tan b loss process of HBA36 IA32 HQ.32; (-) normalised DMA stiffness of

polyarylate; ( ... ) tan b loss process of polyarylate.


Given that raising the Tg in crystalline polymers can lead to increasingly high Tm and intractability, one potential route around this flaw in MCLCPs is to design high-Tg, non-crystalline MCLCPs and it has been found that incorporating units such as 3,3'-diaminodiphenyl sulphone into the MCLCP V can raise the Tg to over 200°C.8 ! There is, however, no modulus retained above the Tg•

MCLCPs with higher levels of crystallinity, such as polymers based on VI, behave like polymers III and IV in the temperature regime 0-120°C but above 120°C exhibit excellent retention of modulus

{o-@-gl t°-@--@-iJ~-@-t " 2" 2"

compared to these polymers.83 This improved temperature performance, however, is at the expense of processability, as the relatively high melting temperature of these polymers (T m > 400°C) renders them difficult to process on conventional injection-moulding equipment.

Up to this stage in this section, the relationship between changing the polymer composition and how this effects the molecular morphology has been discussed. For a given formulation, however, subtle changes can also be observed by controlling the way in which the component monomers are put together-i.e. random versus blocky versus regular alternating copolyesters. For example, for the series of polymers based on the units shown in polymer VI, the alternating copolyester based on (-HBA-BP-HBA-T A-)n- is more crystalline and intractable than the random co polyester analogue.84

Recently, Stupp et al.8s 87 have shown with lower-melting MCLCPs based on structures VII and VIII that the copolymer with the random structure VII melts out of its liquid crystal phase gradually

tg-"<[email protected]"""t~~}-{())--oP-@~J '"

432 W.A. MacDonald

II II II II II II {

0 x"'O;~"'O; x+y",~; x+y-l =3 0 0 O}

0-@-0-C-@----0-C-(CH2)5-C-0-@-C-0-@-0-C-(CH2)5-C n

whereas the regular alternating copolymer VIII exhibits a sudden change, despite the copolymers having identical average lengths of chains. The implications of these observations for bulk mechanical properties are not clear, but studies of this type should contribute significantly towards an understanding of the state of matter of MCLCPs.

8.5 PROPERTIES OF THERMOTROPIC MAIN CHAIN LCPs

8.5.1 Rheology The rheology of MCLCPs is complex. In general, the shear viscosity of MCLCPs is much lower than that of conventional polymers at a comparable molecular weight, and the transition from the isotropic state to the liquid-crystalline state is generally accompanied by a significant decrease in melt viscosity. At the onset of nematic behaviour, the melt viscosity of the MCLCPs is three decades less in order of magnitude than that of a similar but non-mesogenic polymer. 13

The unsheared melt is assumed to have a polydomain texture as shown in Fig. 8.14. 88 Initially shear thinning at low shear rates corresponds to a reduction in domain size. The domains tumble in the flow, the rheology

Z' 'iii o u If)

:;

Stress or Ti me

Fig. 8.14. Relationship between morphology and rheology.


is paste-like and no net long-range orientation is induced. Further reduction in domain size produces an increased surface area and a higher-viscosity state is obtained. At higher shear rates a second shearthinning regime is entered, domain size reduction reaches a lower limit, and if the applied stress is sufficient domain coalescence and long-range orientation can be achieved.

Molecular orientation occurs readily during melt flow and it has been demonstrated that elongational/extensional flow is primarily responsible for orientation of MCLCPs during melt processing.61 ,89 Elongational flow is introduced by drawing-down during fiber spinning and is also present in the fountain flow at the advancing front of an injectionmoulding process. MCLCP melts have very long relaxation times,90 retaining this orientation after flow has ceased and producing anisotropic articles. 91

From a commercial point of view the low melt viscosity of MCLCPs at high shear rates relative to conventional polymers is one of the key features 92 of MCLCPs and enables their use in the following ways:

(a) Injection moulding of component with long or complex flow paths and thin sections.

(b) At very high filler loadings (up to 70% by weight). (c) As a processing aid (bulk lubricant) for conventional thermoplas

tics (providing the processing temperatures of the two polymers overlap). In general the addition of 10% w/w of MCLCP to a conventional thermoplastic approximately halves the melt viscosity of the polymer so that it becomes twice as easy to process.93 This drop is viscosity is greater than that observed with isotropic immiscible blends and the reason for its occurrence is not fully understood.

8.5.2 Mechanical Properties Mechanical properties, particularly tensile strengths and stiffness, depend upon the degree of orientation achieved. This is limited to some extent by the fabrication method and type of article produced, as shown schematically in Fig. 8.15. Thus, a compression-moulded un oriented LCP has mechanical properties similar to that of a conventional isotropic polymer. On injection moulding, tensile bars of MCLCPs generally show superior mechanical moduli to that of conventional glass-fibre-reinforced isotropic thermoplastic (Fig. 8.16),

434

3000

2000

1000

Tensile Strength in Flow Direction MPa

100

50

W.A. MacDonald

//

/~ .// Melt Spun

./ Fibre 5-10!,m // diameter

(J Extruded Film and Tape <1 mm thick

/J Injection Mouldings V <50mm thick

I • Isotropic

Unoriented Solid

5 10 20 30 40 50 60 70

Tensile/Flexural Modulus GPa

Fig. 8.15. Schematic illustration of how tensile strength and stiffness depend upon the degree of orientation achieved.

16

14

12 a;-Il.

~ 10 III :I "3 8 1:J 0 ::!:

~ 6 :I X III 4 u::

2

0 LCP PES PEEK NYLON 66 PST PPS

Fig. 8.16. Flexural modulus of thermoplastics.92 ~, unfilled; D, glass fibre filled.


with the highly oriented skin making the major contribution to the stiffness of the moulding. 92 As the level of elongational flow increases, the mechanical properties increase to the very high mechanical moduli and tensile strengths demonstrated by MCLCP fibers. 14•94

As a simplification, the layered structure of an injection moulding can be considered to be a microcomposite made up of layers in which the direction of reinforcement changes from layer to layer. The observed mechanical modulus of the overall moulding will reflect the integrated effect of the component layers. The important factors governing the modulus would be expected to be the thickness of the component layers and the direction and degree of orientation of the polymer chains within the layers. These factors will largely depend upon the rheological conditions and on the intrinsic response of the polymer chains of the flow conditions. Variations in the mechanical stiffness can also be achieved by altering the formulation. The phase diagram of the polymer series based on polymer IV indicates a change from isotropic to liquid-crystalline melt at around 20% HBA. 78,95,96 Studies on the flexural modulus of injectionmoulded tensile bars over the isotropic to liquid-crystalline range 15% < HBA < 36% showed a smooth systematic increase in modulus rather than a pronounced increase in modulus as the polymer formed an LC mesophase with increasing content of HBA (Fig. 8.17).78 Three selected

15

13

11 • E t7

'" Z 9 <!l

'" :::J 7 :;

'tl 0 :::;

~ 5

:::J )( Q)

u::: 3

0 0 10 20 30 40

% HBA

Fig. 8.17. Flexural modulus versus composition for the series of formulations based on polymer IV.78

436 WA. MacDonald

examples of moulded bars from this series were examined by X-ray diffraction to quantify the chain orientation function as a function of depth of the mouldings. 96 Although skin--core morphology was observed in these mouldings, it was shown that the increase in modulus observed across the series was mainly a consequence of the increase in molecular orientation within the layers. The conclusion was drawn that in order to design a system with high self-reinforcing modulus, it is not sufficient just to ensure that the polymer forms a liquid-crystalline mesophase. It is also necessary that the polymer mesophase is capable of becoming highly oriented along the flow field so that there is a high global orientation, especially within the critical skin regions; i.e. the molecular linearity of the chain in addition to the rheological effects discussed in Section 8.5.1 can have a pronounced effect on bulk mechanical properties.

MCLCPs are tough materials and a benign failure is experienced on impact, similar to that exhibited by long-fibre-reinforced polymers or natural wood; i.e. the failure is neither ductile nor brittle and the mouldings generally do not shatter.

As discussed in Section 8.4.2.3 MCLCPs exhibit a significant fall-off in modulus with temperature and, although this can be influenced by control of the component parts of the MCLCP, this is obviously undesirable in an engineering resin. However, this is partly offset by the exceptionally high stiffness and strength of MCLCPs at room temperature and useful mechanical properties are retained at high temperatures (>200"C).

Fabricated MCLCP articles are anisotropic and the anisotropy ratio, i.e. the difference in properties along and across the flow direction, increases with the degree of orientation and so is highest in fibres. Injection mouldings exhibit anisotropy ratios of between 4: 1 to 10: 1 depending upon mould thickness. 97 Anisotropy increases with decreasing thickness as the proportion of skin to core increases.7. 88 Interestingly, whereas the introduction of fillers tends to increase the anisotropy ratio of conventional isotropic thermoplastics, the introduction of fillers to LCPs disrupts the alignment of the LCP molecules and reduces the anisotropy ratio.97 This is illustrated in Fig. 8.18 (unpublished results from our laboratory) and has the benefit of improving the properties in the across-flow direction.

One of the major problem areas in injection moulding of MCLCPs is the poor weld-line strength exhibited by mouldings because the flow fronts do not knit easily together. This is illustrated in Fig. 8.19, where it

6

5

4

Anisotropy 3 ratio (0'/90')

2


0+------,------,-------,------,------,------, o 10 20 30 40 50 60

Glass content (%)

Fig. 8.18. Comparison of the effect of glass fillers on the anisotropy of a typical MCLCP (SRP) and on polyethersulphone (PES) and polypropylene (PP).

can be seen that the double-gated tensile bar has approximately 10% of the tensile strength of the single-gated moulding. The incorporation of fillers can improve the situation, but not significantly. The problem is currently being overcome by taking careful control over the mould design to move weld lines into areas where they least effect properties. However, the weld-line problem is an inherent flaw of MCLCPs and may prove to be a major 'Achilles Heel'.

MCLCPs absorb very low levels of moisture (typically less than 0·2% on immersion in water) and therefore the change in dimension of mouldings of MCLCPs due to moisture absorption is very low. The coefficients of linear thermal expansion of MCLCPs are much lower than those for conventional polymers (even when glass-fibre reinforced) and are comparable to those for metals, as shown in Fig. 8.20.92 This similarity in thermal expansion for metals and MCLCPs is expected to result in good component integrity and minimal strain when components containing metals (e.g. solder) and MCLCPs are in contact and are subjected to thermal cycling or shock. With conventional

Ga

te

No

We

ld L

ine

Te

ns

ile

10

0

Str

en

gth

1%

) 90

80

70

60

50

40

30

20

10

0

30

% G

lass

_

Fib

re F

ille

d

....L

__

-1. _

_ ~I-

Un

fill

ed

No

W

eld

W

eld

Un

e U

ne

Fig

. 8.

19.

Str

engt

h of

an

MC

LC

P a

t a

wel

d lin

e.

Ga

te .-W

eld

Lin

e

Ga

te

.j::>

U

J

00

~

A s:: ~ '6

§: '" is:

()

~ c. .s "E Q)

·0

~ o ()

c: o ·iii c: as c. ~ OJ E ~ .:: I-


90~--------------------------~,,--------.

80

70

60

50

40

30

20

10

Copper Solder PES NYLON 66 PPS Aluminium LCP PEEK PST

Fig. 8.20. Linear thermal expansion coefficient of thermoplastics.92 D, glass fibre filled; Sl, unfilled.

plastics such thermal treatment can cause separation of the metal and plastic parts of the component owing to differential expansion.

MCLCPs also exhibit very low mould shrinkage and minimal sinkage and warpage compared to conventional isotropic polymers (Fig. 8.21),92

1.7

1.5

1.3

~ 1.1 Q)

Ol as "" 0.9 c: ~ CI) 0.7 "0 :; 0 0.5 ~

0.3

0.1 0

LCP PES PEEK NYLON 66 PST PPS

Fig. 8.21. Mould shrinkage of thermoplastics.92 D, glass fibre filled; Sl, unfilled.

440 W.A. MacDonald

permlttmg precision moulding of components. This is largely related to the unusual molecular morphology of MCLCPs, with which there is very little change in the general configuration of the molecules before and after melting. With isotropic crystalline thermoplastics, on the other hand, the formation of chain-folded lamellae can give rise to warpage/shrinkage problems. However, as mentioned in Section 8.4.2.3, as 'kinks' are introduced into the backbone, the polymer backbone gains more flexibility and crystals of relatively high perfection are observed. This may have an adverse effect on the low warpage/ shrinkage associated with MCLCPs, especially on annealing. The level of kinks in the MCLCP backbone must therefore be taken into consideration when designing MCLCPs for precision moulding applications.

Thus, the presence of flexibility in the MCLCP backbone, in addition to affecting the molecular linearity, may also have a significant effect on warpage and shrinkage by altering the crystallization behaviour of MCLCPs. These factors must be taken into account when designing MCLCPs for high-performance precision moulding applications.

8.5.3 Miscellaneous Properties Gas transport studies on the MCLCP based on polymer III have shown that this polymer has excellent barrier properties, and at 35°C the permeability coefficient for He, Hz, 0z, Ar, N z and COz in this polymer are comparable to or smaller than those for polyacrylonitrile, which is one of the least-permeable polymers known.98 This low permeability seems to stem from the low solubility of these gases in the MCLCP rather than from low transport mobility. The extent of conventional crystallinity for this LCP would have to be 90% or more to explain these solubility results, assuming that the part that is not crystalline in the conventional sense has 'normal' gas solubilities as for glassy polymers such as polyacrylonitrile. This is not the case, as can be seen from Table 8.4, and it therefore seems likely that the liquid-crystalline order is responsible for the very low gas solubilities observed. Significantly, initial unpublished studies carried out in our laboratories on MCLCPs containing 'kinks' of the type illustrated by polymer IV indicate that although the barrier properties are still good they are not as good as for the rod-like polymers based on frustrated packing discussed above. This is presumably related to the ability of the polymer chains to pack and again indicates that the inherent rod-like nature of the polymer backbone can have a significant effect on bulk properties.


MCLCPs have excellent resistance to a wide range of organic solvents and exhibit very good hydrolysis resistance, and the retention of properties in both acidic and basic environments is very good. 14 As with gas barrier properties it is likely that the liquid-crystalline order is responsible for this excellent chemical resistance, and again it is found that the chemical resistance of MCLCPs decrease if kinks are introduced into the backbone.14

8.6 APPLICATIONS

The differences in behaviour between MCLCPs and conventional isotropic polymers and the unique set of properties resulting from the behaviour of MCLCPs are summarized in Fig. 8.22. The main application for MCLCPs will be in areas that exploit combinations of the key properties such as strength, easy flow, excellent dimensional stability, the ability to incorporate high levels of fillers and excellent chemical resistance. Examples of these can be seen from applications identified in the electronics industry such as high-precision parallel interconnecting devices introduced by ITT Canon Inc. for use with large dot-matrix, highresolution, flat-panel displays in such applications as desktop and portable computers, word processors and data-entry terminals. These devices are not permanently attached to the displays, as their driver boards and the alignment of these connectors with the respective termination points

Conventional polymer Self Reinforcing Polymer

I ~~S \ ~!~: t

\ 1.. ') ~ \APPIY shear (~\l~O' (e.g InJecllOn ~ mould)

c.''''''''1i~~1 l, '---'--c~'--'-,.--'-J

Melt

I~~~I\ Ordered

Solid

• Low melt viscosity/ Easy processability High levels of fillers

Very low mould shrinkage

Very low warpage

Anisotropic properties

High modulus (i.e. self reinforcing)

Good solvent resistance

Low water absorption

Excellent barrier properties

Fig. 8.22. Differences in behaviour between MCLCPs and isotropic polymers and summary of the MCLCP key properties.

442 W.A. MacDonald

is critical. The low coefficient of thermal expansion of MCLCPs allows the highly demanding alignment conditions to be met. Further examples are surface-mounted connectors which exploit the excellent dimensional stability after vapour-phase soldering at 215°C coupled with the low thermal expansion which avoids stress on the soldered joints when the assembly is thermally cycled. Other applications are in fibre-optic coupling discs, components for compact disc, bobbins in microwave ovens, automotive uses and packing for chemical towers; and MCLCPs are increasingly being used in demanding applications where metals, ceramics, thermosets and other high-performance polymers were previously employed.

One of the major brakes on the expansion of the MCLCP market has been the relatively high cost of the polymers resulting from expensive feedstocks, and considerable research activity in industry has centred on attempting to lower monomer costs. In addition to this there has also been considerable effort blending MCLCPs with other thermoplastics with the objective of combining the optimum properties of both components while producing the most cost-effective material for a given application. An example of this is a MCLCPjnylon blend which exhibits a very low coefficient of thermal expansion.99 The blends area is likely to be an area of major market activity for MCLCPs in the future.

8.7 CONCLUSIONS

The exceptional balance of properties, combined with the ability to tailor the properties of MCLCPs for particular end uses, has created new opportunities for designers and materials specifiers and these specialists are discovering that they can create new applications. Simplistically, the key properties that are generating so much interest in industry as outlined in Fig. 8.22 can be directly related to the unique molecular morphology of MCLCPs, which in turn controls the micro- and macromorphology. The molecular morphology is in turn directly related to the polymer architecture and it has been shown in this chapter that changing the polymer composition, and to a lesser extent the order in which the monomers units are put together, can have a significant direct effect on the properties of the fabricated articles. The complex interrelationship between polymer architecture, hierarchy of morphology and bulk properties is currently attracting more attention,


and molecular modelling studies currently being carried out in several laboratories should throw considerable light on the inherent rod-like nature of MCLCPs of different formulation and the interrelationship with downstream properties. A deeper understanding of this interrelationship is fundamental to fine-tuning and extending the properties of this unique class of polymers.

ACKNOWLEDGEMENTS

The author would like to thank Drs D.J. Blundell, I.S. Miles, R.A. Chivers, A.H. Windle and A.M. Donald for helpful comments in the preparation of this chapter.

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10. Pletcher, T.c., U.S. Patent 3991013 and 3991014, 1976 (E.I. DuPont de Nemours and Co.).

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Ward. Elsevier Applied Science, London, 1987. 16. Zhou, Q.F. & Lenz, R.W., Polymer Prepr., 1983, 24(2), 255. 17. Zhun, Q., Lenz, R. & Jin, J.-I., in Polymeric Liquid Crystals, ed. A. Blumstein.

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444 W.A. MacDonald

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20. Jin, J.I., Antoun, S. & Lenz, R.w., Br. Polymer J., 1980, 12, 132. 21. Strzelecki, L. & Van Luyen, D., Eur. Polymer J., 1980, 16,301 22. Ober, CK., Jin, 1.1. & Lenz, R.W., Adv. Polymer Sci., 1984, 59, 101 21 Ciferri, A., in Polymer Liquid Crystals, ed. A. Ciferri, W.R. Krigbaum & R.B.

Meyer, Academic Press, New York, 1982. 24. Roviello, A. & Sirigu, A., Macromol. Chern., 1982, 183. 895. 25. Blumstein, A. & Blumstein, R.B., in Recent Advances in Liquid Crystalline

Polymers, ed. L.L. Chapoy. Elsevier Applied Science, London, 1985. 26. Griffin, A.C & Havens, SJ., J. Polymer Phys., 1981, 19,951. 27. Abe, A., Macromolecules, 1984, 17, 2280. 28. Aguilera, C, Bartulin, 1., Hisgen, B. & Ringsdorf. H., Makromol. Chern., 1983,

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31. Jackson, WJ. & Kuhfuss, H.F., J. Appl. Polymer Sci., 1980,25, 1685. 32. Demartino, R.N., J. Appl. Polymer Sci., 1983, 28, 1805. 33. Griffin, B.P. & MacDonald, W.A., European Patent W5527, 1980 (Imperial

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European Patent. 303935, 1987 (Bayer A.G.). 35. Pielartzik, H., Ebert, W., Meyer, R.V., Traenckner, H.J. & Ostlinning, E.,

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U.S. Patent. 4668760, 1984 (Owens-Corning Fibreglassl. 40. Kuhfuss, H.F. & Jackson, W.J., U.S. Patent 3778410, 1973 (Eastman Kodak

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Polymeric Liquid Crystals, ed. A. Blumstein. Plenum Press, New York, 1985. 71. Mitchell, G.R. & Windle, A.H., Colloid & Polymer Sci., 1985,263, 230. 72. Biswas, A. & Blackwell, J., Macromolecules, 1988,21, 3152. 73. Hanna, S. & Windle, A.H., Polymer, 1988,29, 223. 74. Windle, A.H., Viney, c., Golombok, R., Donald, A.M. & Mitchell, G.R.,

Faraday Discuss. Chern. Soc., 1985, 79. 75. Golombok, R., Hanna, S. & Windle, A.H., Mol. Cryst. Liq. Cryst. 1988, 155,

281. 76. Spontak, R. & Windle, A.H., Polymer, 1990,31, 1395. 77. Lemmon, TJ., Hanna, S. & Windle, A.H., Polymer, 1989, 30, 2. 78. Blundell, DJ., Chivers, R.A. & MacDonald, W.A., High Pe~f'ormal1ce Poly

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446 W.A. MacDonald

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Chapoy. Elsevier Applied Science, London, 1985. 89. Viola, G.G., Baird, D.G. & Wilkes, G.L., Polymer Eng. Sci., 1985,25,888. 90. Ide, Y. & Chung, T.S., J. Macromol. Sci., 1984/85, B23, 497. 91. Huynh-ba, G., ACS Polymeric Mater. Sci. Eng., 1984,50, 141. 92. Cox, M.K., Mol. Cryst. Liq. Cryst. 1987, 153,415. 93. Cogswell, F.N., Griffin, B.P. & Rose, J.B., European Patent 30417,1981. 94. Prevorsek, D.e., in Polymer Liquid Crystals, ed. A. Ciferri, W.R. Krigbaum &

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1699. 99. Chung, T.-S., Plastics Eng., 1987,43(10),39.

Chapter 9

Applications of lCP Materials

Jan-Fredrik Jansson Polymeric Materials, Royal Institute of Technology, Stockholm, Sweden

9.1 INTRODUCTION

The industrial development of thermotropic liquid crystal polymer (LCP) materials can be traced from its theoretical origins, through the identification of useful compositions, to full commercialization. The future industrial challenge will be to define and develop applications which take advantage of the unique properties of these materials.

Representative classes are found in wholly aromatic polyesters, aromatic-aliphatic polyesters, wholly aromatic polyesteramides, aromatic-aliphatic polyesteramides, aromatic polyazomethines, aromatic polyester-carbonates, etc.

Thermotropic LCP materials have very good thermal, physical, dielectric, optical and mechanical properties as well as good chemical resistance, low flammability, very good dimensional stability, etc. In addition to their good product properties they show remarkable ease of processing thanks to their low melt viscosity and high melt strength.

Already commercial demands are growing in the fibre-optic, chemical process, electric/electronic, aerospace, automotive and household markets and new areas of potential applications continue to emerge.

Starting from the classical liquid-crystalline copolyester (PET /PHBA) developed by Eastman Kodak and first reported to be liquid-crystalline

447

448 J an-F redrik Jansson

in 1974, the Carborundum terephthalic acid/4,4' -biphenoijPHBA allaromatic co polyester which was reported to be injection mouldable in the same year, and the Du Pont LCPs based on substituted hydroquinones from 1975, there are a number of commercial LCPs available today: Xydar (Dartco, 1984)*, Vectra (Hoechst/Celanese, 1985), Ultrax LCP (BASF), Grannlar (Granmont), Rodrum LC-5000 (Unitika) resins. Du Pont and Bayer have also recently introduced new amorphous and crystalline LCP grades on the market. All the LCPs are also produced in highly filled variants. Table 9.1 shows some of the milestones in the development of LCPs up to 1985 as described by Calundann. 1

Table 9.1 Milestones in the industrial development of thermotropic polymer liquid crystal

line materials

1940-1956

1956

1965

1972

1974

1975-1976

1975

1982-

1982-

1983-

1984

1985

1985

1985

Lyotropic biopolymers - Tobacco mosaic virus - Collagen - Poly(y-benzyl-L-glutamate)

Theory-Flory - 'Rigid rod': anisotropic solution

Lyotropic aromatic polyamides - Kevlar

Melt-processable wholly aromatic polyester - Ekkcel 1-2000 (Carborundum)

Thermotropic pol/ester - Aromatic/aliphatic copolymers Z7G (Eastman)

Wholly aromatic polyester, polyazomethines

Intense patent activity - Industrial and academic interest increasing

Teijin, Rh6ne-Poulenc, and others

Monsanto test marketing Thermotropic polyaryelate (BPA)

ICI test marketing Celanese LCPs

Dartco Mfg. (Dart-Kraft) Commercial with Xydar

Celanese commercial with Vectra

Sumitomo to be commercial with 'E Konol' fibre

Mitsubishi Chern., Unitika test marketing 'X-7G' variants

* Now available from Amoco.

Applications of LCP Materials 449

It has also been found that the properties of ordinary engineering polymers can be greatly improved by blending them with thermotropic liquid crystal polymers.2- 4 Usually the processing properties are also improved. The fact that the materials are biphasic is not a disadvantage. The predominantly flexible phase contains some liquid-crystalline sequences which provide a certain amount of reinforcement in these regions. Costs are lower than for pure liquid-crystalline materials.

The thermotropic materials may be processed by injection moulding, extrusion, thermoforming, blow moulding, etc. To modify their properties they are usually filled with glass fibres, or carbon fibres or other inorganic fillers.

Some liquid-crystalline materials contain smectic phases which give ferroelectric and pyroelectric properties and are proposed for use in optical shutters, displays and heat sensors, as has been discussed by Simmonds in Chapter 7.

High-quality LCP coatings have been developed for non-bake applications. Some of them can also be cross-linked.

Jackson,S Chung, Calundann and East6 and others have given extensive reviews of the development, applications and future trends of LCP technology. Table 9.2 summarizes some applications pointed out by the Hoechst-Celanese Corporation. Hoechst-Celanese estimates the market for thermotropic LCPs in the mid-1990s to be 73% consumer, 8% industrial, 5% electrical/electronics, 5% telecommunications, 4% transportation, 3% aircraft/aerospace and 2% other.7

Table 9.2 Areas of application of liquid crystal polymers (Hoechst/ Celanese)

Electronics/electrical

Fibre-optics

Automotive

Industrial

Chemical process

Domestic equipment

Other

Connectors, surface-mount components, relays bob-bins, capacitor housings, potentiometers, switches

Strength members, couplers, connectors

Fuel-system components, electrical systems

Motor components, lamp housings, conveyor belt components, gears

Tower packings, pump housings, pump shafts, valves

Compact disc components, microwave equipment and turntables

Medical components, watch components, safety equipment, chemical analysis equipment, leisure goods

450 Jan-Fredrik Jansson

Thermoplastic liquid crystal polymers generally have lower viscosities, longer melt relaxation times and higher melt strengths than ordinary thermoplastics. Little or no extrudate swelling occurs despite the very high melt elasticity. Thermotropic liquid crystal polymers can be processed in ordinary thermoplastic processing equipment and, because of the low viscosity and low extrudate swelling, even more easily than ordinary thermoplastics.

9.2 INJECTION-MOULDED PRODUCTS

Liquid crystal polymer materials offer many advantages over conventional thermoplastics in injection moulding, including low mould shrinkage, minimum warpage and distortion, fast cycle time, ability to mould thin parts, low moisture absorption, and excellent chemical resistance. The mechanical properties are comparable to those of materials filled with short glass fibres.

Although relatively few LCP materials have been commercialized, a large number of liquid crystal polymers suitable for injection moulding are described in the literature and in patents. Generally, LCP grades containing fillers or additives are chosen with improved overall properties.

The important advantages are somewhat counteracted, however, by the difficulty of controlling the orientation and anisotropy of the material and the low strength of the weld lines. Usually a multilayer structure is observed in injection-moulded parts, including a highly oriented skin, a very slightly oriented or completely non-oriented core, and several intermediate layers with differently oriented structures.813 Obviously the conditions during both the filling and the packing stage of the moulding cycle play a more important role in the formation of this multilayer structure than in the corresponding phenomena in ordinary thermoplastics.

Several methods have been developed for reducing anisotropy. One way is to increase the mould temperature, but this obviously also increases the costs. Multiaxial processing is a new technique in which the biaxial flow involving lubricated squeezing between discs has shown some advantages. 14- 17 Manipulation of the moulding geometry and the local flow has turned out to be very effective.1s IBM has developed a technique involving moulding through a rotating dye to introduce both radial and circumferential orientation. 19 21


Push-pull moulding developed by Brunel University, England, and Klockner Ferromatik Dessma in Germany and others offers an interesting way of reducing anisotropy and also strengthening weld lines. 22 The idea is to fill the cavity from different directions using two injection units, a microprocessor and a mould with double melt channels or hot runners and two separate gates. During the injection cycle the melt flow from the two units alternates until the freezing of the gates, introducing a repeated switching of the flow direction and a series of layers from skin to core having different orientation directions.

Blending and addition of fillers are other ways of reducing anisotropy. Thus, 50% glass fibre or other high-modulus mineral fibre reduces the skin-core effect and widens the processing window without changing the mechanical properties substantially. A large number of LCP grades suitable for injection moulding have been developed. Hoechst-Celanese has presented a long series of exclusive injection moulded products. 2 3

LCP materials have shown excellent performance in surfacemounted components, i.e. mounting of electronic components directly onto the surface of a printed circuit board instead of using holes in the board. The application requires materials having high mechanical strength and modulus, durability and toughness, chemical resistance, flame retardance, suitable coefficient of thermal expansion, dimensional stability and ease of processing. The LCP material also exhibits less dimensional change during vapour-phase soldering than, for example, PBT, PPS (glass-fibre filled) and PEL Thus, highpincount, micro-miniature and precision/high-duty connectors, pin grid arrays and burn-in sockets are among many applications being evaluated. Figure 9.1 shows a miniature surface mounted connector, and Fig. 9.2 demonstrates a high-precision parallel interconnecting unit with 768 contacts.

A series of advanced rotary switches has been designed based on good impact resistance, heat resistance to soldering temperature, good insulation properties, solvent resistance and flame retardance.

LCP materials offer several advantages over metals and traditional thermoplastics for use in compact disc scanning heads (Fig. 9.3). They maintain their high stiffness even in narrow sections (wall thickness down to 0·3 mm), and retain their specific modulus and other mechanical properties even at the elevated temperatures generated by a laser.


Fig. 9.1. Surface-mount miniature connector made from Vectra (by courtesy of Hoechst-UK Ltd, Polymers Division).

Fig. 9.2. The ITT Cannon 'Parallel Interconnect' made from Vectra (by courtesy of Hoechst-UK Ltd, Polymers Division).

Optical couplers are used to split, combine or distribute light beams to and from optical fibres. The material of the coupler must have a coefficient of thermal expansion close to that of the optical fibres, excellent dimensional stability at elevated temperatures, high modulus, and extremely low mould shrinkage to allow very narrow tolerances. A special grade of glass-fibre-filled LCP has been developed for this application. 24

Cost savings of 50% to the end user have been achieved by replacing the nickel-plated zinc die-cast diode receptacle for fibre optics with LCP material. Critical dimensions must be held within a 0·009 mm tolerance, as indicated in Fig. 9.4.


Fig. 9.3. Lens holder for compact disc scanning made from Vectra (by courtesy of Hoechst-UK Ltd, Polymers Division).

A number of advantages have been reached in using graphite flake modified LCPs in saddle packings in distillation columns instead of the ceramic packings. The packings were tested with great success in a plant for the production of formic acid.

Rivets made from LCP materials offer many good qualities compared to metal fasteners. LCP rivets used to fasten carbon fibre/epoxy composites are cheaper and lighter than, for example, titanium fasteners and given fewer corrosion problems than aluminium. LCP rivets are also favourable in connection with aluminium panels.

Small gear vessels in car windshields, washer pumps, etc., have been made from LCP materials containing 50% glass and mineral fillers and a lubricant. Very tight moulding tolerances are required in combination with dimensional stability and good friction properties.


Fig. 9.4. Diode receptacles manufactured from Vectra (by courtesy of HoechstUK Ltd, Polymers Division).

The phenolic/glass-fibre compound conventionally used in a telecommunication bobbin was replaced by LCP and this gave better processability and higher heat resistance.

A bakery pan for use in magnetic conveyor systems has been produced by injection-moulded LCP including integrated ferrous or magnetic material members. 25

Speaker vibration members with excellent acoustic properties have been manufactured from a LCP material modified by glass and carbon fibres, wollastonite, talc, mica and glass flakes. 26

9.3 EXTRUDED LCP RODS AND PROFILES

Large-diameter rods extruded from LCP materials are used to replace steel wire.6 Many advantages are also gained in using LCP rods as


strength members in optical keyboard applications. LCP rods have low weight, considerable flexibility and excellent tensile properties which make handling easier and prevent the optical fibres from breaking during the lay-down process.

Like many oriented structures, LCP rods have very small or even negative coefficients of thermal expansion, which reduce stresses caused by temperature variations. Other favourable properties are their good chemical resistance and almost zero water uptake.

Thermotropic LCP rods are manufactured by extrusion under conditions which produce a strong molecular orientation in the flow direction. The orientation is preserved by rapid cooling to freeze the material in a time shorter than its orientational relaxation time. However, the slow heat transfer from the fluid to the surrounding cooling medium may cause a temperature distribution within the rod which gives the internal region time to relax and lose its orientation. Particularly in thick rods, the modulus may be considerably lower in the middle than in the surface regions. To achieve a high order of orientation and high modulus it is necessary to choose a material which has a high degree of molecular rigidity and packability and is easy to crystallize. It has been shown that copolyesters containing p-HBA or 6-hydroxy-2-naphthonic acid as a main component are able to orient to high levels and exhibit relatively high moduli in rods having diameters of 1-3 mm.37 It has been discovered that some of the LCP copolymers have superior moduli and load-bearing properties equal to those of pultruded epoxy/glass fibre rods.

Melt-extruded members are also used as support for optical fibre cables,27 and are produced in different geometries to give different types of support for the optical fibre bundles. The polymers used have a modulus in the range of 60 GPa, a tensile strength of 700 MPa and an elongation at break of 1·5% in the longitudinal direction.

9.4 ORIENTED SHEETS AND FILMS

Several methods have been developed for the manufacture of films and sheets from thermotropic LCP materials. Bayer has presented a technique to form oriented sheets and films from fibres by pressing at a temperature at which the fibres melt to form a matrix without losing their molecular orientation.28 For unidirectional mouldings moduli of the order of 50-60 GPa at a tensile strength of 400-500 MPa were obtained


in the fibre direction, corresponding to more than 90% of the fibre properties.

In a process developed by Hoechst-Celanese, a biaxially oriented film is produced29 by extruding an anisotropic dope consisting of a polymerization solution of a LCP and a suitable solvent. This film is oriented transversely to the extrusion direction to enhance the transverse strength, after which the film is solidified and the solvent removed.

In addition to their good mechanical properties, LCP films have very low coefficients of diffusion which make them interesting as highperformance packaging materials.

9.5 THERMOFORMING AND BLOW MOULDING

Sheets of mineral-filled LCP materials have been used for thermoforming and electroplating for printed circuit boards. 6 During the fabrication, calendering rolls have to be placed close to the die. The stiffness of the sheets requires a spool with a diameter about 200 times the thickness of the product.

Biaxially oriented products have been made by extrusion blow moulding using a die temperature below the melting temperature. However, the biaxial orientation as well as the properties vary through the thickness of the product and depend strongly on the thermomechanical conditions of the process. For LCP materials with slow crystallization rates, blow moulding in the supercooled state represents a possible method of obtaining multidimensional properties. 30.31

9.6 MATRIX MATERIALS FOR COMPOSITES

LCP/carbon-fibre composites have been developed and are used as secondary composites within the aerospace industry.6.32 LCP matrices offer low viscosity for the impregnation of the fibres, and excellent chemical resistance. Impregnation is achieved in a cross-head die, after which the composite panels are prepared by compression moulding the stacked prepreg sheets.

The mechanical properties of the LCP matrix/carbon-fibre composites are comparable to those of conventional epoxy/carbon-fibre materials. The flexural modulus retention at high temperatures is very good, whereas the flexural strength decreases, owing to the poor bonding


between the fibres and the matrix. The impact resistance is superior to that of conventional thermoset systems.

9.7 FIBRES

High-performance polymeric fibres have the potential to revolutionize the use of composites and other fibre applications. Although at present only aramid fibres and high-modulus polyethylene fibres have achieved a truly commercial breakthrough, it can be expected that a large number of fibres based on thermotropic LCPs will be available in the near future, such as Twaron (Enka), Technora (Teijin), Ekonol (Sumitomo) and Vectran (Celanese-Kuraray).3

Although the properties differ somewhat between different types of polymeric fibres, they have many positive qualities in common: low density, very high specific tensile properties, high toughness, good chemical resistance. Negative qualities include high anisotropy, low compressive and shear strengths and negative axial coefficients of thermal expansion. The application temperature covers a broad range from that of low-melting polyethylene to a useful range of more than 300°C in some compositions, but it is generally much lower than for conventional fibres. Creep may also be a problem with some compositions.

The moduli of polymeric fibres are determined mainly by the molecular orientation, but depend also on the inherent stiffness of the molecules, the intermolecular interaction, the crystallinity and the packing density of the molecules, while the chemical composition defines the temperature range within which the fibre can be used.

All LCP fibres have an extremely high chain orientation with a Hermans orientation number greater than 0·95. Usually the modulus lies between 100 and 200 GPa with a tensile strength of more than 2 GPa.

The solution spinning of lyotropic LCPs such as Kevlar and Twaron (both poly(p-phenylene terephthalamide)) and poly(1,4-benzamide) is well documented. Thermotropic LCPs can be converted into fibres using melt spinning facilities able to produce suitable elongational flow fields. The oriented structure of the fibres appears in several hierarchical levels.2

Table 9.3 summarizes the different stiffness classes of polymeric fibres and their formation. 34

The mechanical properties develop quickly with increasing draw ratio and are further enhanced by heat treatment at temperatures slightly

458 Jan·Fredrik Jansson

Table 9.3 Formation of high·performance polymeric fibres

Relative chain stiffness

Flexible, linear 'conventional' polymer

Aromatic or aromatic heterocyclic 'con· ventional' polymer

'Semi·stiff' mesogenic polymers

'Rod·like' mesogenic polymers

Examples

Polyethylene, polyacrylonitrile

Hydrazide·aramid copolymers, poly· oxidiazoles

Thermotropic copoly· ester, lyotropic aramid

Fibre formation processes

Superdrawing, gel spinning

Solution spinning, drawing

Melt spinning, dry·jet wet spinning, heat treament

Poly(p·phenylenebenzo· Dry jet-wet spinning, bisthiazole) heat treament

below the melting point in an inert environment for minutes or hours. The heat treatment usually also improves the chemical resistance and thermal properties. Both the tensile strength and the creep properties improve with increasing molecular weight, although a very high molecular weight may lead to problems in forming the desired highly oriented structure.

The unique properties of the lyotropic LCP fibres make them suitable for a number of applications. Generally available are Kevlar (Ou Pont) and Twaxon (Akzo) poly(p·phenylene terephthalamide) (PPT). PPT is used in ropes for helicopter winches, racing sheets, halyards, etc. where high specific modulus and tensile strength are required. Ribbon parachutes made from PPT have only half the weight of the corresponding polyamide design. Fabrics made from LCP fibres are used as garments in helmets and military jackets. These types of helmets are said to be able to stop a flying bullet in a very short distance. Figure 9.5 shows the effect of a bullet on the material. The excellent tear resistance and thermal insulation properties also make PPT suitable for gloves and ciothings for hostile environments.

PPT fibres are used in advanced composite applications in combination with epoxy. However, to attain good properties the fibres have to be pre-dried to a very low moisture content. To overcome the weak compressive and shear strength of the PPT fibres in applications where these properties are crucial, hybrids with carbon fibre can be used. High-performance composites made from laminated PPT and biaxiallY oriented polyester films are used in racing sails,


Fig. 9.5. Twaron (Akzo) fabric panel after a shooting test with a 9-mm parabellum. Missile and deformed bullet nose are shown for comparison (reproduced by

courtesy of the Institute of Metals)38

in some applications together with high-modulus polyethylene fibres. Fishing boats, motor yachts, canoes and pressure vessels have been manufactured from PPT composites and show outstanding performance.

PPT fibres are used in combination with other types of fibres and fillers in brake linings to replace asbestos. The strength and toughness and the dimensional and thermal stability of the fibres contributes to the low creep and cracking resistance of these new materials.

Chung, Calundann and East give an extensive review of the formation, structure and application of LCP fibres. 6 Some suggested applications are summarized in Table 9.4.


Table 9.4 Applications of LCP fibres

Sample

Protective fabrics Ballistic vests Gloves Clothing

Strong fabrics Tarpaulins Conveyer belts

Coated fabrics Inflatable boats Sails

Industrial fibres Ropes, cables (oil rig mooring) Ropes (marine uses) Filament wound pressure vessels Sails Sewing threads

Rubber reinforcement Radial tyres (trucks, automobiles)

Belts (conveyor, power transmission)

Plastics reinforcement Space applications

Aircraft, interior

Aircraft, propeller

Boats, canoes, kayaks

Military (helmets) Sporting goods

Cement reinforcement Building material~, pipes

Friction uses, asbestos replacements Brake linings, clutch facings, gaskets packing

Property

Strength, thermal resistance Cut, thermal resistance Strength, thermal resistance, comfort

Strength, weatherability Strength, stiffness, fatigue resistance

Strength, fatigue resistance Strength, fatigue resistance

Strength, low density, low creep Strength, low density, low creep Strength, low density, low creep Strength, weatherability, low fatigue Strength, weatherability, low fatigue

Strength, stiffness, low fatigue, adhesion, low density, oxidation resistance

Strength, stiffness, hydrolytic stability

Tensile properties, low density, thermal resistance

Tensile properties, low density, thermal resistance

Tensile properties, low density, thermal resistance, vibration damping

Tensile properties, low density, low fatigue, vibration damping

Tensile properties, low density Modulus, compression strength

Tensile properties, hydrolytic stability, thermal properties

Modulus, non-toxicity, non-aggressive wear, thermal stability


9.8 COATINGS

Reduction of solvent level in coatings is important from both an ecological and an economic point of view. However, lowering the amount of solvent raises the viscosity, which has to be compensated for by lowering the molecular weight and/or the glass transition temperature of the binder. The general conditions and applications of LCP binders have been discussed by Pappas. 3 5

A low molecular weight or a low glass transition temperature may affect the final properties of the coating film and tends to narrow the application window, i.e. the range of bake temperature, the time needed for the coating film to reach acceptable properties, and so on. Thus, a low molecular weight or a low glass transition temperature requires binders with higher functionality, lower equivalent weight and a higher concentration of reactive groups to build up the desired final properties of the coating. Polymers with high functionality and reactive groups tend also, however, to have large intermolecular interactions which increase the viscosity and the surface tension. Low molecular weight also increases the risk of sagging during spray-up on vertical surfaces and during baking of the coating.

The design of resin systems for high-solids coatings based on epoxides, urethanes and thermosetting acrylics and polyesters, air-dry alkyd coatings, etc, is a particularly challenging area. The important characteristics of these systems also include the dry-to-touch and through-cure times. The desired low viscosity is reached by effective blending of compatible polymers such as low- and high-glass-temperature polyesters, acrylics and polyesters or urethanes. Conventionally, dry-totouch is achieved by solvent evaporation while curing involves oxidative cross-linking. Often dual-cure systems are used in which the solvent evaporation is complemented by oxidative cross-linking as the dryto-touch mechanism and the through-cure is based on other cross-linking processes.

It has been found that the addition of different kinds of LCPs to coating resins gives several advantages. LCP binders provide favourable combinations of hardness and flexibility, together with good adhesion and fast drying.

Thus, liquid crystallinity in alkyd resins provides several important practical benefits. Solution viscosity is reduced by the formation of nonaqueous dispersions, drying times are sharply reduced and films are both


hardened and toughened. LC acrylic lacquers also have two other major advantages as binders for higher-solid non-bake coatings: they form concentrated, stable, low-viscosity dispersions in common solvents and the films have an extraordinary combination of high hardness and high impact resistance.

REFERENCES

1. Calundann, G.W., in High Performance Polymers: Their Origin and Development, ed. R.B. Seymour & G.S. Kirshenbaur, Elsevier, London, 1986.

2. Brostow, W., Polymer, 1990, 31, 979-995. 3. Brostow, W., Dziemianowicz, T.S., Romanski, 1. & Werber, W., Polymer Eng.

Sci., 1988, 28(12), 785-795. 4. Dutta, D., Fruitwala, H., Kohli, A. & Weiss, R.A., Polymer Eng. Sci., 1990,

30(17),1005-1018. 5. Jackson, W.J., Jr, Mol. Cryst. Liq. Cryst., 1989, 169,23-49. 6. Chung, T.-S., Calundann, G.W. & East, A., Handbook of Polymer Science and

Technology, ed. N.P. Cheremisinoff, Marcel Dekker, New York, 1989, vol. 2. 7. Dole, 1.R. Chemtech., 1987, 17, 242. 8. Ophir, Z. & Ide, Y, Polymer Eng. Sci., 1983, 14, 792. 9. Hedmark, P.G., Rego Lopez, M., Westdahl, M., Werner, P.-E., Jansson, 1.-F.

& Gedde, U.W. Polymer Eng. Sci., 1988,28, 1248. 10. Jackson, W.J., Jr. & Kuhfuss, H.F., J. Polymer Sci. Phys., 1976, 14, 2043. 11. Engberg, K., Knutsson, A., Werner, P.-E. & Gedde, U.W., Polymer Eng. Sci.,

1990, 30(24), 1620--1627. 12. Gedde, u.-W. & Engberg, K., personal communication. 13. Duska, 1.1., Plastics Eng., 1986, 12, 39. 14. Chartraei, S., Macosco, e.W. & Winter, H.H., J. Rheol., 1981,25,433. 15. Soskey, P.R. & Winter, H.H., J. Rheol., 1985, 29, 493. 16. Lin, YG., Winter, H.H. & Lieser, G., Liq. Cryst., 1988,3, 519. 17. Lin, YG. & Winter, H.H., Liq. Cryst., 1988,3,593. 18. U.S. Patent, 4.843.109, Jun. 27, 1989. Shaped articles formed from polymers

capable of exhibiting anisotropic melts. (Imperial Chemical Industries, pic). 19. Zachariades, A.E. & Economy, 1., Polymer Eng. Sci., 1983,23,266. 20. Economy, 1., Volksen, W. & Geiss, R.H., Mol. Cryst. Liq. Cryst., 1984, 105,

289. 21. Economy, 1., J. Macromol. Sci. Chem., 1984, A21, 1705. 22. Miller, B., Plastics World, Nov. 1990, No. 25. 23. Hoechst Celanese; Plastics in Engineering, Issues, 34, 3:(1986), 1:(1987),

2(1987), 1:(1988). 24. Chen, J.e., Gorky, D.V., Haley, R.e., Jaarsma, F.e. & McChesney, e.E.,

Proc. 5th Speretec, Ohio Univ., 1985, p. 61. 25. U.S. Patent 4.922.811, May 8, 1990. Bred pan fabricated of liquid-crystal

polymer.


26. U.S. Patent 4.880.591, Nov. 14, 1989. Method for manufacturing speaker vibration member.

27. U.S. Patent 4.910.057, Mar. 20, 1990. Melt extruded elongated member suitable for improved service as a stiffening support in an optical fibre cable.

28. U.S. Patent 4.923.660, May 8, 1990. Process for the production of mouldings and films from thermotropic polymer.

29. U.S. Patent 4.898.924, Feb. 6, 1990. Process for the production of biaxially oriented rigid rod hetero-cyclic liquid crystalline polymer films.

30. Blizard, K.G. & Baird, D.G., Int. Polymer Processing, 1989, IV(3) 172-178. 31. Blizard, K.G., Wilson, T.S. & Baird, D.G., Int. Polymer Processing, 1990, V(l). 32. Chung, T.S. & McMahon, P.E., J. Appl. Polymer Sci., 1986,31,965. 33. Jaffe, M., Calundann, G. & Yoon, H.-N., Handbook of Fiber Science and

Technology: Vol. Ill. High Technology Fibres. Part B., ed. M. Lewin & 1. Preston· Marcel Dekker, New York, 1989.

34. Calundann, G., Jaffe, M., Jones, R.S. & Yoon, H., Fibre Reinforcements for Composite Materials, A.R. Bunsell, Elsevier, London, 1988.

35. Pappas, S.P., J. Coatings Technol., 1989,61(774), 51-53. 36. Chen, D.-S. & Jones, F.N., J. Appl. Polymer. Sci., 1989,37, 1063-1078. 37. Itoyama, K., J. Polymer Sci, Part C. Polymer Lett., 1989,27, 369-375. 38. Collyer, A.A., Mater. Sci. Technol., 1990,6,981-992. 39. U.S. Patent 4.803.235. Feb. 7, 1989. Composition for injection moulding.

(Poly. Plastic Co. Japan). 40. Jackson, W.J., Jr, J. Appl. Polymer Sci. Appl. Polymer Symp., 1985, 41,

25-33.

Index

Acrylate-methacrylate-chlorocrylate sequence, 221

Acrylate-water system, 266 Activation energy, 220, 221 Addition polymerization of meso genic

monomers, 384-7 Alkenes, hydrosilylation of, 390, 391 Alkyl polyoxyethylene surfactants, 267 a-helical polypepetides, 6 a-process, 181, 186, 199-202, 204, 215,

217, 219, 222, 223, 228 aI-process, 216, 218, 219 a2-process, 216, 218, 219 af3-process, 181, 183, 184, 202, 215 Amino acid residue, 265 Amphiphilic molecules, 240, 241, 253 Amphiphilic monomers, 255 Anisotropic potential, 281, 283 Anisotropy

molecular shape, of, 10, 31, 32 physical properties, of, 162 polarizability, of, 284 reducing, 450, 451

Anisotropy factor, 8 Aramid fibers, 296 Argand diagram, 148 Argand plot, 163 Aromatic polyamides, 274-6 Asymmetric diffusion, 155 Autocorrelation function, 147

Backbone, 359-65 chain, 189 constitution of, 360

Back bone-contd. dynamics, 212-23

Banded textures, 12 Barotropic materials, 9 Batonnets, 62, 63, 65 Benzazole polymers, 276 Benzene pyridine ring, 374 Benzothiazole polymers, 276 Benzyl ether linkage (OCH2), 376 f3-process, 181, 182, 183, 184, 186, 199,

201, 202, 204, 209, 223 13" -process, 224 f3l-process, 208, 209, 211 13 2 -process, 208, 210, 211 f3-relaxation, 133, 135 Biforked compounds, 39 Biological membranes, 267 Biparallel structures, 20 Bisphenols, 413 Bis-swallow-tailed compound, 40 Block and graft PLCs, 14 Blow moulding, 456 Body-centred cubic structure, 247 Bowlic molecules, 38 Bulk polymerization, 385

13C NMR, 106, 107, 117, 118, 121, 126, 127, 130-2, 136, 139

Calami tic compounds, 34 Calamitic mesogens, 371 Calami tic phases, 22 Canonic mesophases, 35, 158 Cationic polymerization, 387

465

466 Index

Central mesogen linkages, 375-7 Chain formation, 296-8 Chain scission, 340 Chain stiffness, 220 Chiral C* phase, 228 Chiral combs with heterocyclic

mesogens, 397 Chiral compounds, 50, 60 Chiral meso gens, 356 Chiral nematics, 160 Chiral smectic C mesophases, 395 Chiral smectic C phase, 50 Chiral smectic C* phase, ferroelectric

modes in, 171-3 Chiral smectogens, 357 Chlorohydroquinone, 413 Cholesteric homopolymers, 61 Cholesteric LCP devices, 393 Cholesteric liquid crystals, 12 Cholesteric nematics, 160 Cholesteric phases, 11, 356, 357 Cholesteric textures, 60-2 Cholesterics, 48 Cholesteryl ester copolymers, 378 Cholesteryl esters, 367 Chromatographic applications, 400,

401 Class rx, 15 Class p, 15 Class y, 15 Class s, 19

subclass sD, 19 subclass eO, 19 subclass sP, 19

Class K, 19 Class )., 20

subclass ).1, 20 subclass n, 20 subclass le3, 20

Class cp, 19 Class 1/1, 20

subclass 1/11, 20 subclass 1/12, 20

Class 8, 20 subclass 81, 20 subclass 82, 20

Class ( subclass (M, 18

C1ass-contd. subclass (R, 18 subclass (S, 18

Clearing point, 283 Clearing temperature, 13 Coagulation process, 310-15 Coatings, 449, 461, 462 Cole-Cole function, 149 Cole-Cole parameter, 218 Cole-Cole plots, 148, 149, 152 Cole-Cole relationship, 163 Columnar phases, 263 Columnar structure, 22 Comb LCPs, 19, 349, 351

see also Thermotropic side-chain liquid crystal polymers

Comb nematics, 88 Combined architectures, 357 Compact disc scanning heads, 451 Complex dielectric constant, 148 Complex dielectric permittivity, 146,

147, 151 Composites, matrix materials for, 456,

457 Compressive strength of fibers, 334-7 Condensation polymerization, 388 Condis crystals, 3 Cone-shaped molecules, 38 Conformation of LCPs, 87-91 Conformational disordering, 3 Conformational isomerism, 376 Conic molecules, 21 Contact method, 71 Contour projection length, 282 Coordination polymerization, 365 Copoly (n-butyl, nonyl-isocyanate)

(PBNIC), 192, 193 Copolymerization, 408, 410 Copolymers, 377-83

dynamic mechanical analyses, 429 mesogens with non-mesogens,

377-9 with two different mesogens,

379-81 Creep, 324-9 Creep-rupture lifetimes, 339 Critical micelle concentration

(CMC), 241, 242, 254, 261, 264

Index 467

Cross-linked polymers, 358, 381-3, 389 Crystalline liquids, 31 Crystalline smectics, 51, 52, 77 Cubic phase, 246, 247

bicontinuous, 249 Cyano-group, 224 Cyanophenylbenzoatoxy side group,

218 Cybotactic nematic meso phases, 83, 84 Cybotactic nematics, 47 Cylindrical micelles, 264

Davidson-Cole function, 149 Debye equation, 147, 163 Degree of orientation, 295 Degree of polymerization (DP), 360,

362 15 loss peak, 216 15-process, 208, 213, 219, 220, 222, 227 15'-process, 216, 222, 15"-process, 216 15-relaxation, 133 Deuteron NMR spectra, 89 Dialkyl surfactants, 240 4,4-diamiodiphenyl, 273 Dielectric constants, 353 Dielectric loss peak, 215, 217 Dielectric loss spectrum, 203 Dielectric loss versus frequency and

temperature, 229 Dielectric permittivity tensor, 163 Dielectric relaxation, 7

macromolecular liquid crystals, in, 143-236

polymers, in, 174-86 uniaxial phase, in, 163-71

Dielectric relaxation spectra, 163 Dielectric relaxation spectroscopy,

principles of, 144-56 Dielectric spectroscopy of liq uid

crystals, 156-73, 187-231 Dielectric susceptibility, 145 Differential scanning calorimetry

(DSC), 7, 337, 352 Differential thermal analysis (DT A),

337 4, 4' -di-n-heptyloxyazoxybenzene, 13

Dimethyl formamide (DMF), 390 Diode receptacles, 452 Dipole autocorrelation function, 154,

156, 177 Dipole correlation function, 152, 153,

183 Director, 8, 11, 159, 188, 215, 226

alignment parallel with, 350 alignment perpendicular to, 350 orientation of, 160, 161

Disc-{;omb structures, 19 Disc-like amphiphiles, 264 Disc-like mesogens, 32 Disc-like micelles, 243 Disc-like molecules, 34, 36 Disc-like structure, 158 Discogens, 37, 38 Discoidal amphiphiles, 253 Discotic meso gens, 358 Discotic polymers, 18, 204 Discotic systems, 357-9 Discotic thermotropic polymers, 122,

123 Dissado-Hill function, 150 Distillation columns, 453 n-dodecyl hexaethylene glycol

monoether, 241 Dodecyl hexaethylene glycol

monoether (C12E06j-water system, 249

Double (or combined) PLCs, 20 Double systems, 357-9 Draw ratio, 290, 291, 293-5 Dynamic mechanical analysis curves,

329 Dynamic mechanical measurements

(DMA),328 Dynamic mechanical testing (DMT), 7

Elastic behavior, 318-24 Elastic constants, 353 Electron diffraction studies, 298-305 Electron microscopy, 305-10 Electronic components, 451 Electro-optical applications, 391-9 Electro-optical behavior, 7 Enantiotropic materials, 15

468 Index

End-over-end reorientations, 166, 168, 171, 184, 227

Euler angles, 154 Excluded-volume theories, 281 Extruded LCP rods and profiles, 454,

455

Ferroelectric LCPs, 229-31, 393-7 Ferroelectric modes in chiral smectic

C* phase, 171-3 Ferroelectric properties, 449 Ferroelectric siloxane combs, 396 Fiber composites, 4 Fiber optics, 452 Fibers, 457-9

applications, 460 compressive strength of, 334-7 creep, 324-9 formation, 293 mechanical properties, 315-37 morphology, 296-315 strength of, 330-4 stress relaxation, 324-9

Fillers, 451 Films, 455

mechanical properties, 315-37 morphology, 296-315 strength of, 330-4

Flexible chains, 178 Flexible polymers, 5

in bulk, 179-84 Flexible spacers, 365-70, 376, 411, 412

influence of length of, 85 role of, 89

Flip-flop motion, 217 Flip-flop reorientation, 216, 228 Flory-Huggins theory, 74 Flow behavior, 286-90 Focal-conic domain, 62 Free-radical polymerization, 364 Free-volume concept, 171 Frequency-dependent dielectric

constant, 146 Frustrated chain packing, 410, 411 Fuoss-Kirkwood curves, 216 Fuoss-Kirkwood equation, 149

y-process, 208, 211 y-relaxation, 133 Gas-liquid (partition)

chromatography, 400, 401 Gaussian hypergeometric function, 150 Glass-liquid transition, 123-5, 175 Glassy liquid crystal, 60 Goldstone mode, 172, 173, 230 Grandjean steps, 61 Graphite flake modified LCPs, 453

HI phase, 260 1 H NMR, 105, 125-7 2H NMR, 105, 106, 118, 126, 128-30,

133, 135, 136, 139 Havriliak-Negami function, 149 HBA/HNA copolymers, 198 HBA/PET copolymers, 198, 199 Hemiphasmidic compound, 40 Heterogeneous composites, 4 Hexagonal phases, 246, 264 Hexatic smectics, 50, 51 Hierarchical phase structures, 7 High-melting polymers, 417 Highly polar end groups, 53-5 Homeotropic texture, 59, 60 Homopoly (n-hexyl isocyanate)

(PHlC), 192, 193 H2S-AgNO) penetration technique, 306 Hydrophilic head group, 239 Hydrophobic effect, 239 Hydrophobic tail, 239, 240 Hydrosilylation of alkenes, 390, 391 4-hydroxybenzoic acid, 418

Injection moulding, 450-4 In-situ composites, 7 Intelligent processing, 7 Internal friction, 8 Inverse combs, 19 Ionic polymerization, 387 Isotropic diffusion, 156 Isotropic phase, 156, 168, 189, 219 Isotropic state, 212-23

Index 469

Kerr effect measurements, 186 Kevlar, 296 Kohlrausch-Williams-Watts (KWW)

function, 149, 150, 183

Laboratory frame of reference (LF), 153, 161, 165

Ladder polymers, 276-8 Lamellar phases, 246, 248, 251, 264 Larmor frequency, 106 Lateral substituent, 373-4 Lath-shaped mesogens, 356 LC phases, 263 Lecithin, 240 Legendre polynomials, 161-2 Lennard-lones potential, 283 Liposomes, 269 Liquid crystal elastomers, 399-400 Liquid crystal mesophase, 157 Liquid crystal polymers (LCPs), 1, 7

advantages over metals and traditional thermoplastics, 451

applications, 25, 447-63 areas of application, 449 blends, 4, 7 classification, 14-21 conformation of, 87-91 dielectric spectroscopy of, 187-231 first reported, 188 history, 2 mechanical properties, 4 milestones in development of, 448 molecular structures, 15 networks, 21 polymorphism, 55-67 properties of, 24-5 structural modifications, 42 textures, 55-67

Liquid crystal twins, 14 Liquid-crystalline materials, clas-

sification and terminology, 1 Liquid-crystalline molecules, 163, 169 Liquid-crystalline order, 195 Liquid-crystalline solid glass, 195 Liquid crystallinity, 1-30, 32

experimental work, 11 nature of, 8-11

Liquid crystals, 3, 31, 157, 169 dielectric spectroscopy of 156-73 phases of, 11-14 rigidity of, 4

LMMLCs, binary mixtures of, 73 Local dynamics

longitudinal liquid crystal polymers, in, 127-32

mesomorphic polymers, in, 123-36 side-chain thermotropic polymers,

in, 133-6 Local field factor, 151 Longitudinal liquid crystal polymers,

15 local dynamics in, 127-32

Longitudinal thermotropic polymers, NMR studies of orientational and conformational order, 107-15

Loss angle, 146 Loss factor, 146 Low-angle X-ray diffraction, 250, 262 Low-molar-mass (LMM) systems, 32 Low-molar-mass liquid crystals

(LMMLCs), 34-41, 73 Lyotropic liquid crystalline polymers,

187 Lyotropic liquid crystallinity, 187 Lyotropic liquid crystals, diffraction

pattern, 250 Lyotropic main-chain liquid crystal

polymers, 273-348 flow behavior, 286-90 mechanical properties, 315-37 order in, 278-96 spinning process, 290-6 synthetic aspects, 273 thermal properties, 337-40

Lyotropic materials, 9, 32 Lyotropic one-comb amphiphiles, 10 Lyotropic polymers, 191-4 Lyotropic side-chain polymer liquid

crystals, 237-72 phase behavior, 256-67 synthesis, 253-6

Lyotropics, 157

Macromolecular liquid crystals,

470 Index

Macromolecular liquid crystals -contd. dielectric relaxation in, 143-236

Macroscopic correlation function, 147 Macroscopic polarization

autocorrelation function, 148 Magic-angle spinning (MAS), 106 Magnetic field effects on SCLCPs, 89 Magnetic susceptibility, 8 Maier-Saupe potential, 283 Maier-Saupe theory, 74, 278, 281-5 Main-chain LCPs (MCLCPs), 32,41-

4, 188, 190, 195, 197-204, 223-8 chain sequence extension, 87 enhanced ordering in, 87 methods for reducing transition

temperatures, 43 orientational order in, 87

Marbled texture, 56 Mark-Houwink coefficient, 419 Mark-Houwink exponent, 278 Mark-Houwink relation, 280 Matrix materials for composites,

456,457 Mechanical properties, 315-37 MEPSIL solvents, 400 Mesogen length, 371 Mesogen linkage, 365 Mesogen phase correlations, 379 Mesogenic monomers, addition

polymerization of, 384-7 Mesogenic units, 189, 190,212-23, 226 Mesogens, 31, 32, 356, 357, 370-7

lateral attachment of, 367-70 Mesomorphic phases, 31, 212-23, 223,

227 Mesomorphic polymers

local dynamics in, 123-36 slow motions in, 137, 138

Mesophases, 157 characterization, 31-101, 352 stability, 361 types, 3

meta-aromatic isomers, 414 N-methylglucamide surfactants, 260 Methyloxyphenylbenzoatoxy side

group, 218 Microscopic correlation function, 151,

169

Miscibility tests, 70-6 Mixed structures, 20 Moire fringes, 60 Molecular composites, 4

aerospace industry, 6, 7 Molecular dipole moment, 165 Molecular dynamics, 175, 176, 178, 179 Molecular frame of reference (mF), 153,

161 Molecular mass, 360-2 Molecular properties, 21-3 Molecular structures, 15, 21-3, 32-45 Monodomains, 161, 163 Monomer liquid crystals (MLCs), 1

anatomy of, 14 history, 2

Monosaccharide amphiphiles, 260 Monotropic materials, 15 Monotropic mesophases, 352 Multiaxial processing, 450 M ultidomain, 164

Nematic domain, 59 Nematic droplets, 56 Nematic-isotropic transition

temperature, 283 Nematic liquid crystal, 11 Nematic marbled texture, 59 Nematic mesophases, 158 Nematic phases, 13, 46, 47, 158, 164,

167, 188, 189, 219, 226, 292 Nematic systems, theory of, 10 Nematic textures, 55-60, 251 Nematic-to-nematic transitions, 370 Nematogenic mesogens, 356 Network LC elastomers, 381-3 NMR studies, 105-7

principles of, 125-7 thermotropic polymers, 103-41

Non-aqueous dispersion polymerization, 418

Non-linear links, 412-16 Non-linear optical (NLO) effects, 398,

399 Non-polar SCLCPs, 79 Normalized autocorrelation functions,

155, 165

Index 471

Novel nematic mesophase, 78

Optical couplers, 452 Optical data storage, 392, 393 Optical microscopy, 305-10 Order parameter, 8, 162, 163,283, 285 Orientability parameter, 293 Orientation parameter, 321 Orientational disordering, 3 Orientational order parameters, 104 Orientational polarization, 146, 147 Orientational probability distribution

function, 154 Oriented mono layers, 267-70 Oriented nematics, 82 Oriented samples, X-ray diffraction

patterns, 82-6 Oriented sheets and films, 455, 456 Orienting potential, 282 Orthogonal liquid crystal polymers, 15 Orthogonal phases, 78

Packing constraints model, 244 Para-aromatic polyamides, 285, 296,

305, 315 Parallel (biparallel) systems, 367-70 Parallel structures, 20 Paramorphic fan-shaped texture, 67 Partial rigidity, 10 Partially fixed SCPs, 76 PBLG solutions, 186 PBNlC, 192, 193 PBO, 273, 286, 287, 310, 315, 318, 329,

332, 333, 335 PBT, 6, 273, 276, 286, 287, 298, 315,

317, 318, 329, 332-35, 337, 339 PBZ, 276, 278 PET, 198, 427 PET 6OA20B20, 201 PET,A(l_,)2B(l_,)2' 201 PET,Ay/2By/2 , 202 PETjHBA,_" 200 PET4Q/HBA60 , 199 PET/PHBA,447 Phase diagrams, 7, 71, 74, 76, 191, 249,

260, 311

Phase structures, 7, 21-3 Phase transitions, 7, 353, 361, 366, 368,

369, 372, 374, 375, 395 Phase-transfer catalysis (PTC), 389,

390 Phasm-like mesogens, 32 Phasmidic compound, 39 p-phenylenediamine, 273 Phenyl/naphthyl copolyesters, 202 PHIC, 191, 192, 193 Photochemical radical polymerization,

385-7 Photochromic LCPs, 401 Photo imaging technology, 401 Piezoelectric elastomers, 400 Piezoelectric properties, 402 Plastic crystalline phase, 157 Plastic crystals, 3, 157, 160 PLCs, 1, 7 PmPTA,323 Polar forked compound, 39 Polar groups, 176 Polar side groups, 81 Polarizing optical microscopy, 262 Polyacrylate-based side-chain liquid

crystals, 206 Polyacrylates, 115-18, 133-6,217,

262, 263 Polyacrylic liquid crystals, 210 Polyacrylics, 220 Poly(n-)alkyl isocyanate) (PAlC)

polymers, 179, 186, 192, 194 Polyamide-hydrazides, 273 Poly(I,4-benzamide) (PpBA), 279-81 Poly(p-benzamide) (PpGA), 273 Poly(4,4-benzanilidylene

terephthalamide, see PpBA T Poly(y-benzyl-L-glutamate), 179, 273 Poly(n-butylisocyanate) (PBIC), 179 Polycatenar compounds, 38-41 Polydispersity, 360-2 Poly(ester amides) (PEAs), 11 Polyflexibility, 14 Poly(n-hexyl isocyanate) (PHlC), 191 Poly(p-hydroxybenzoic acid)

(PHB), 6 Polymer chains, 176, 177, 179, 180 Polymer effect, 356, 359

472 Index

Polymer liquid crystals, see Liquid crystal polymers

Polymer membranes, 268 Polymer modification reactions, 389,

390 Polymerization, 253, 254, 259, 261, 262,

265, 267-70 363-5 Polymers

dielectric relaxation in, 174-86 diluted solutions, in, 177-9

Poly(meth)acrylamide, 264 Polymethacrylates, 115-18, 133-6 Polymethyacrylate, 262, 263 Poly(p-oxybenzoyl-co-ethylene

terephthalate) (X7G), 417 Polyphenyl, 6 Poly(p-phenylenebenzobisthiazole), see

PBT Poly(p-phenylene terephthalamide)

(PPT), 6, 458, 459 Poly(1,4-phenylene terephthalamide),

see PpPTA Poly(p-phenylene terephthalamide), see

PpPTA Poly(phenylhydroquinone-co

terephthalic acid), 420 Polyphosphoric acid (PPA), 277, 278 Polysiloxane-based side-chain LCPs,

207, 208 Polysiloxanes, 118-21, 136, 206, 216 Positional disordering, 3 Post-shearing relaxation, 12 Powder samples, X-ray diffraction

patterns, 77-81 Power-law relations, 282 PpBA, 274, 283, 293, 296, 297, 299, 310,

312,315,323 PpBA T, 283, 296, 297, 300, 339 PPD,276 PpPT A, 273, 279-83, 286, 287, 291,

292,296-8,300,301,305,306,308, 310-14, 316, 318, 321-6, 328-34, 337-9

PpTA,274 PRD-49, 296, 339 Pre-draw ratio, 294, 295 Proportionality factor, 145, 147 Proton dipolar decoupling, 106

Push-pull moulding, 451 Pyramid-like mesogens, 32 Pyramidic molecules, 38 Pyroelectric properties, 449

Quaternary ammonium (meth-)acrylates, 254

Realignment phenomena, 372 Re-entrant nematic (RN) phases, 14 Refractive indices, 305 Relaxation processes, 133 Reorientation by instantaneous jumps,

153 Reversed cubic phase, 246 Reversed hexagonal phase, 246 Rigid biphenyl moiety, 259 Rigid polymers, 6 Rigid rod heterocyclic (ladder)

polymers, 276-8 Rigidity of liquid crystals, 4 Rivets, 453 Rod-like mesogens, 32 Rod-like molecules, 34-6, 45-55 Rod-like polymers, 184-6, 191 Rod-like structure, 158, 179 Rod-shaped micelles, 243 Rotary switches, 451 Rotational diffusion, 153 Rotational dynamics, 196 Rotational mobility, 157 Rotational molecular dynamics, 153,

154

Saddle packings, 453 Sanidic packing, 22 SANS, 89, 90 Schiff's bases, 71 Schlieren texture, 56, 58, 67, 251 Schroeder-van Laar equation, 72 Secondary transitions, 123-5 Semi-crystalline liquid-crystalline solid,

195 Semi-flexibility, 10 Semi-flexible polymers,S, 6

Index 473

Semi-rigid polymers, 187 Shear flow-induced birefringence, 288 Sheets, 455 Siamese twin mesogen, 14 Side-chain dipole motions in glassy

state, 209-12 Side-chain LCPs, 32, 188-91, 196,

205-28 magnetic field effects on, 89 mixtures of, 74 polymer backbone of, 88 with and without flexible space, 349

Side-chain liquid crystalline polymers (SCLCPs), 44, 45

Side-chain mesogens, 190 Side-chain thermotropic polymers

local dynamics in, 133-6 NMR studies of orientational and

conformational order, 115-21 Si-H infrared adsorption, 255 Siloxane copolymers with malonate

diester mesogens, 378 Simple parallel structures, 20 Single crystals, 161 Slow motions in mesomorphic

polymers, 137-8 Small-angle neutron scattering (SANS),

88-90 Small-angle scattering methods, 87-91 Small-angle X-ray scattering (SAXS),

89, 301, 304 Small-molecule liquid crystals

(SMLCs), 349 Small-molecule systems, 377 Smectic A fluctuations, 47 Smectic A phases, 49, 53, 62, 66, 67 Smectic C fluctuations, 47 Smectic C phases, 50, 53, 54, 65-7, 83 Smectic C* phase, 228 Smectic E phases, 81 Smectic F phase, 50 Smectic phases, 12, 13,48-50, 158, 160,

188, 189, 219, 226, 449 Smectic textures, 62-7 Sodium dodecyl sulphate (SDS), 245 Sodium octanoate, 251 Sodium lO-undecenoate, 254 Soft mode, 172, 173

Solid phase, 156, 175 Solid polymers, dielectric relaxation

in, 181 Spacer groups, 350 Spacer length, 365-7 Spacer linkage, 355 Spherulites, 251 Spinning process, 290-6, 311

air-gap, 292 SQUID magnetometer, 8 Star LCPs, 15, 253, 358 Stiff chains, 178, 179, 278 Stochastic stationary Markov process,

154 Stress-optical effects, 399 Stress relaxation, 324-9 Structure formation, 310-15 Supramolecular meso gens, 22 Surface-mounted components, 451 Surface-stabilized ferroelectric liquid

crystal (SSFLC), 394 Surfactants, 239-52

classification, 240 dilute micellar solutions, 239-43 ionic, 242 liquid crystal formation of small-

molecule, 245-9 micelle size and shape, 243-5 parts of, 239 penetration experiment, 252 techniques of characterization,

250-2 Surfactant-water system

monomeric, 257 polymeric, 258

Swallow-tailed compound, 40 Symmetric diffusion, 155

Tacticity, 362-5 Tail-to-tail twins, 14 Temperature sensors, 12 Terephthalic acid, 413 Terminal groups, 371, 372 Terminal substituent, 371 Thermal free-radical polymerization,

384, 385 Thermal properties, 337-40

474 Index

Thermally stimulated depolarization, 7 Thermochromic LCPs, 401 Thermoforming, 456 Thermogravimetry (TG), 337 Thermomechanical analysis (TMA), 7 Thermo-optical analysis, 7 Thermotropic LCPs, 41-5, 447-63

processing, 449 Thermotropic liquid crystals, 9, 35, 36 Thermotropic main-chain LCPs,

407-46 anisotropy of, 437 applications, 441-3 backbone flexibility, 440 chain irregularities and chain

stiffness, 425 coefficient of thermal expansion, 442 design, 408-16 entropy of fusion, 425 heat of fusion, 425 higher order containing kinks, 427 history, 407, 408 kinked structure, 426 macromorphology, 421 mechanical properties, 433-40 microstructure, 422-4 miscellaneous properties, 440, 441 molecular morphology, 424-32, 440 morphology, 420-32 mould shrinkage, 439 nature of crystallites, 426 ordered regions, 428 properties, 432-41 rheology, 432, 433 solution and melt characterization,

419,420 strategies for lowering 1m, 414 structure, 420-32 summary of key properties, 441 syn thesis, 416-19

Thermotropic polymers, 32, 157, 195-31

NMR studies of, 103-41

Thermotropic side-chain LCPs, 349-406

applications and applicable materials, 391-402

backbones, 354, 355, 359-65 copolymers, 377-83 general structure features, 353-9 historical development, 351 linkage components, 355 scope and nomenclature, 351-3 structure-property correlations,

359-77 synthesis of comb and parallel

systems, 383-91 Threaded texture, 56, 58 Tractor texture, 12 trans-gauche conformation, 411 Twaron, 296 Twisted nematics, 160 Two-dimensional magic-angle

spinning HMR, 117 Type A polymer, 259

Ultra-high-molecular-weight (UHMW) polyethylene fibers, 335

Uniaxial nematics, 46, 47 Uniaxial symmetry, 159

VI phase, 259, 260 Vesicles, 269

Wedge-shaped samples, 61 Wide-angle X-ray scattering (W AXS),

300, 301

X-ray diffraction, 76-86, 250, 298-305, 311

oriented samples, 82-6 powder samples, 77-81

liquid crystal polymers: from structures to applications

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