Liquid Crystalline Polymers [1st ed.] 9783030433499, 9783030433505

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Polymers and Polymeric Composites: A Reference Series Editor-in-Chief: Kamal K. Kar

Lei Zhu Christopher Y. Li  Editors

Liquid Crystalline Polymers

Polymers and Polymeric Composites: A Reference Series Editor-in-Chief Kamal K. Kar Advanced Nanoengineering Materials Laboratory, Materials Science Programme Indian Institute of Technology Kanpur Kanpur, India Advanced Nanoengineering Materials Laboratory, Department of Mechanical Engineering Indian Institute of Technology Kanpur Kanpur, India

This series provides a comprehensive collection of reference handbooks on all aspects around polymers and polymeric composites. Polymeric materials of all sorts have been emerging as key materials for many applications and for meeting the challenges of the twenty-first century. From commodity applications to engineering and high tech applications, even including aerospace subsystems, these materials have an important role to play. The study of polymeric and polymeric composite materials is one of the most important and of the most vibrant focus areas in chemical and material scientific research. “Polymers and Polymeric Composites: A Reference Series” compiles the most comprehensive reference handbooks on these materials under one roof. Readers will find all they need to know in wellorganized and thoroughly structured reference works covering various topics, such as the structures and properties of polymers, polymeric materials and composites (e.g. structures of amorphous and of crystalline polymers, viscoelastic properties, mechanical and thermal properties, and many more); methods and methodology (including polymer characterization, polymerization reaction engineering, polymer processing, and many more); or different compound classes (from polymer additives, polymer blends, and fiber reinforced composites, to liquid crystalline polymers, nano-polymers and nano-polymeric composites, and even bio-polymeric materials). While each volume is dedicated to a selected topic, concisely structured and thoroughly edited by experts, with contributions written by leading scientists, the complete collection provides the most comprehensive and most complete overview over the entire field of polymers and polymeric composites. Volumes in this series serve as reference compilation for every scientist working with or on polymers, polymeric materials and composites, whether at universities or in industry, from graduate student level to practitioners and lead scientists alike. More information about this series at http://www.springer.com/series/15068

Lei Zhu • Christopher Y. Li Editors

Liquid Crystalline Polymers With 369 Figures and 41 Tables

Editors Lei Zhu Department of Macromolecular Science and Engineering Case Western Reserve University Cleveland, OH, USA

Christopher Y. Li Department of Materials Science and Engineering Drexel University Philadelphia, PA, USA

ISSN 2510-3458 ISSN 2510-3466 (electronic) ISBN 978-3-030-43349-9 ISBN 978-3-030-43350-5 (eBook) ISBN 978-3-030-43351-2 (print and electronic bundle) https://doi.org/10.1007/978-3-030-43350-5 © Springer Nature Switzerland AG 2020 This work is subject to copyright. All rights are reserved by the Publisher, whether the whole or part of the material is concerned, specifically the rights of translation, reprinting, reuse of illustrations, recitation, broadcasting, reproduction on microfilms or in any other physical way, and transmission or information storage and retrieval, electronic adaptation, computer software, or by similar or dissimilar methodology now known or hereafter developed. The use of general descriptive names, registered names, trademarks, service marks, etc. in this publication does not imply, even in the absence of a specific statement, that such names are exempt from the relevant protective laws and regulations and therefore free for general use. The publisher, the authors, and the editors are safe to assume that the advice and information in this book are believed to be true and accurate at the date of publication. Neither the publisher nor the authors or the editors give a warranty, expressed or implied, with respect to the material contained herein or for any errors or omissions that may have been made. The publisher remains neutral with regard to jurisdictional claims in published maps and institutional affiliations. This Springer imprint is published by the registered company Springer Nature Switzerland AG. The registered company address is: Gewerbestrasse 11, 6330 Cham, Switzerland

Preface

Liquid crystalline polymers represent an important class of polymeric materials and have attracted both scientific and technological interests from researchers all over the world. In the past years, this research field has witnessed significant progress in liquid crystalline structure design, synthesis, characterization, theory, and applications. This book is a timely contribution, reporting on recent discoveries in selfassembled structures and dynamics of both rod-like and discotic liquid crystalline polymers and their potential applications in optoelectronics, photoactive actuation, stimuli-response, flame retardancy, gas permeation, and nanocomposites. Chapter 1 presents an introduction of liquid crystalline polymers with different architectures (main-chain versus side-chain) and building blocks (calamitic or discotic). Their self-assembled mesophases include nematic, smectic, and cholesteric. These liquid crystalline polymers find applications in microelectronics, information technology, apparels, and medical devices due to low dielectric loss, mechanical strength, thermal conductivity, and chemical and solvent resistance. From Chapters. 2 to 10, self-assembly and dynamics of liquid crystalline polymers are introduced. For example, Chapter 2 discusses a unique type of side-chain liquid crystalline polymers, that is, mesogen-jacketed liquid crystalline polymers, their synthesis, and molecular weight-dependent columnar phase formation. Chapter 3 introduces liquid crystalline polymers based on discotic building blocks, emphasizing their synthesis, structure, and phase transformation. Chapter 4 discusses anisotropic liquid crystal networks prepared from reactive mesogens and their applications in various optoelectronic devices. Chapter 5 presents a comprehensive description of columnar phase-forming polymers and their complex thermotropic phase transformation. In Chapter 6, liquid crystalline [60]fullerenes are described. These giant molecules can self-assemble into various supramolecular structures because of their unique spherical shapes. Chapter 7 introduces the structure and self-assembly of liquid crystalline block copolymers. In particular, hierarchical structures in this unique system are discussed. Chapter 8 focuses on supramolecular self-assembled discotic liquid crystalline LEGOs. As a result of different covalent attachments of discotic moieties, co-assemblies with intermixed columnar structures are obtained. In Chapter 9, mesoscale structures and dynamics of polymer melts and liquid crystalline fluids are discussed. Intriguingly, when the dimension approaches submillimeter scales, these fluids exhibit finite shear elasticity v

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at low frequency. Chapter 10 presents a comprehensive review of the theory and computer simulations of flow-processes of polymer liquid crystals for lyotropic and thermotropic liquid crystals using the classical Leslie-Ericksen and Landau-de Gennes models. From Chapters 11 to 21, liquid crystalline polymer functions and potential applications are reviewed. Chapters 11 and 12 discuss liquid crystalline conjugated polymers, including their synthesis, structures, and optoelectronic applications. Chapters 13 to 16 focus on photoactive liquid crystalline polymers and block copolymers. In recent years, interests in photoactive alignment and actuation have stimulated active research in this particular area. Chapter 17 introduces a stimuliresponsive liquid crystalline block copolymer based on cyanobiphenyl and poly (ethylene oxide) side chains. Chapter 18 introduces the photoalignment of liquid crystalline molecules using fluorinated polyimides. Chapter 19 reviews the gas permeation and barrier properties of a series of flexible and liquid crystalline vinyl polymers, polyesters, and polyamides. Chapter 20 describes the synthesis and flameretardant properties of high temperature liquid crystalline polymers. Finally, Chapter 21 reviews liquid crystalline polymer nanocomposites. From these contemporary reviews of various aspects of liquid crystalline polymers, this book serves as a reference to the polymer community. In addition, this book provides certain directions for future research opportunities for both academia and industry. August 2020

Lei Zhu Christopher Y. Li

Contents

1

Introduction to Liquid Crystalline Polymers . . . . . . . . . . . . . . . . . Soma Banerjee and Kamal K. Kar

Part I 2

3

1

LCP Self-Assembly and Dynamics . . . . . . . . . . . . . . . . . . . . . .

27

Mesogen-Jacketed Liquid Crystalline Polymers: Molecular Design and Synthesis . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . Zhihao Shen

29

Liquid Crystalline Polymers Derived from Disc-Shaped Molecules . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . Shilpa Setia, Sandeep Kumar, and Santanu Kumar Pal

59

4

Anisotropic Liquid Crystal Networks from Reactive Mesogens . . . Dae-Yoon Kim, Namil Kim, and Kwang-Un Jeong

95

5

Columnar Phase-Forming Polymers . . . . . . . . . . . . . . . . . . . . . . . . Shi Jin

117

6

Fullerene Liquid Crystals . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . Xiaoming Yang, Tiantian Zhu, and Yingfeng Tu

149

7

Structure and Assembly of Liquid Crystalline Block Copolymers . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . Kishore K. Tenneti, Xiaofang Chen, Qiwei Pan, and Christopher Y. Li

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Supramolecular Self-Assembly of Discotic Liquid Crystalline LEGOs . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . Lei Zhu

217

Probing Submillimeter Dynamics to Access Static Shear Elasticity from Polymer Melts to Molecular Fluids . . . . . . . . . . . . Laurence Noirez

249

8

9

10

Liquid Crystalline Polymers: Structure and Dynamics . . . . . . . . . Alejandro D. Rey, Edtson E. Herrera-Valencia, and Oscar F. Aguilar Gutierrez

273

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Contents

Part II

LCP Functions and Applications . . . . . . . . . . . . . . . . . . . . . .

315

11

Liquid Crystalline Conjugated Polymers . . . . . . . . . . . . . . . . . . . . Matti Knaapila, Roman Stepanyan, and Andrew P. Monkman

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Liquid Crystalline Conjugated Polymers with Optoelectronic Functions . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . Kazuo Akagi

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Photodeformable Liquid Crystalline Polymers (LCPs) Lang Qin, Wei Gu, and Yanlei Yu

.........

361

14

Photoactive Liquid Crystalline Polymers . . . . . . . . . . . . . . . . . . . . Asit Baran Samui

391

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Photoresponsive Liquid Crystalline Polymers . . . . . . . . . . . . . . . . Xiao Li and Haifeng Yu

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Light-Sensitive Azobenzene-Containing Liquid Crystalline Polymers . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . Carlos Sánchez-Somolinos

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New Stimuli-Response Liquid Crystalline Polymer Architectures . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . Lalit Mahajan, Dennis Ndaya, Prashant Deshmukh, and Rajeswari Kasi Photoalignment of Liquid Crystal Molecules Using Fluorine-Containing Polyimides . . . . . . . . . . . . . . . . . . . . . . . . . . . Shuichi Sato, Hironaga Matsumoto, Setsuko Matsumoto, and Kazukiyo Nagai Gas Permeation and Barrier Properties of Liquid Crystalline Polymers . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . Shota Ando, Shuichi Sato, and Kazukiyo Nagai

20

Highly Flame-Retardant Liquid Crystalline Polymers . . . . . . . . . . Li Chen and Yu-Zhong Wang

21

Characterizations of Nanocomposites of Liquid Crystalline Polymers . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . Tae Young Ha, Yong-Ho Ahn, Bo-Soo Seo, Donghwan Cho, and Jin-Hae Chang

Index . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . .

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493

523 549

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About the Series Editor-in-Chief

Kamal K. Kar Advanced Nanoengineering Materials Laboratory, Materials Science Programme Indian Institute of Technology Kanpur Kanpur, India Advanced Nanoengineering Materials Laboratory, Department of Mechanical Engineering Indian Institute of Technology Kanpur Kanpur, India Prof. Kar pursued higher studies at the Indian Institute of Technology Kharagpur, India, and Iowa State University, USA, before joining as a Lecturer in the Department of Mechanical Engineering and Interdisciplinary Programme in Materials Science at IIT Kanpur in 2001. He was a BOYSCAST Fellow in the Department of Mechanical Engineering, Massachusetts Institute of Technology, USA, in 2003. Prof. Kar is currently holding the Champa Devi Gangwal Chair Professor of the Institute. Before this, he has also held the Umang Gupta Institute Chair Professor (2015–2018) at IIT Kanpur. He was the Head of the Interdisciplinary Programme in Materials Science from 2011 to 2014 and Founding Chairman of Indian Society for Advancement of Materials and Process Engineering Kanpur Chapter from 2006 to 2011. Prof. Kar is an active researcher in the broad areas of nanostructured carbon materials, nanocomposites, functionally graded materials, nano-polymers, and smart materials for structural, energy, water, and biomedical applications. His research works have been recognized through the office of the Department of Science and ix

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About the Series Editor-in-Chief

Technology, Ministry of Human Resource and Development, National Leather Development Programme, Indian Institute of Technology Kanpur, Defence Research and Development Organisation, Indian Space Research Organization, Department of Atomic Energy, Department of Biotechnology, Council of Scientific and Industrial Research, Aeronautical Development Establishment, Aeronautics Research and Development Board, Defence Materials and Stores Research and Development Establishment, Hindustan Aeronautics Limited Kanpur, Danone Research and Development Department of Beverages Division France, Indian Science Congress Association, Indian National Academy of Engineering, and many more from India. Prof. Kar has more than 220 papers in international refereed journals, 135 conference papers, 10 books on nanomaterials and their nanocomposites, 50 review articles/book chapters, and more than 55 national and international patents to his credits, some of these have over 200 citations. He has guided 18 doctoral students and 80 master’s students so far. Currently, 17 doctoral students, 10 master students, and few visitors are working in his group, Advanced Nanoengineering Materials Laboratory. Dedicated to my wife, Sutapa, and my little daughter, Srishtisudha, for their loving support and patience, and my mother, late Manjubala, and my father, late Khagendranath.

About the Volume Editors

Lei Zhu Department of Macromolecular Science and Engineering Case Western Reserve University Cleveland, OH, USA Professor Lei Zhu received his B.S. degree in Materials Chemistry in 1993 and M.S. degree in Polymer Chemistry and Physics in 1996 from Fudan University. He received his Ph.D. degree in Polymer Science from the University of Akron in 2000. After a 2-year postdoctoral experience at the Maurice Morton Institute, University of Akron, he joined the Institute of Materials Science in the Department of Chemical, Materials and Biomolecular Engineering at the University of Connecticut as an Assistant Professor. In 2007, Lei was promoted to Associate Professor with tenure. In 2009, he moved to Department of Macromolecular Science and Engineering at Case Western Reserve University as an Associate Professor. Lei was promoted to Full Professor in 2013. His research interests include high κ polymer and organic-inorganic hybrid nanomaterials for high energy density capacitor applications, development of artificial antibody as nanomedicines, and supramolecular selfassembly of discotic liquid crystals. Lei a is recipient of NSF Career Award, 3M Non-tenured Faculty Award, DuPont Young Professor Award, and Rogers Teaching Excellence Award. He is author/co-author of 190 refereed journal publications and 9 book chapters. He delivered over 170 invited talks and 45 contributed presentations.

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About the Volume Editors

Christopher Y. Li Department of Materials Science and Engineering Drexel University Philadelphia, PA, USA Christopher Li is a Professor in the Department of Materials Science and Engineering at Drexel University. He received his B.S. degree in Polymer Chemistry from the University of Science and Technology of China in 1995 and his Ph.D. in Polymer Science from the University of Akron in 1999. After working as a postdoc at the Maurice Morton Institute of Polymer Science, University of Akron, for 2 years, he joined the Department of Materials Science and Engineering at Drexel University in 2002 as an Assistant Professor and was promoted to Associate and Full Professor in 2007 and 2011, respectively. His research interests center on the structure and morphology of polymers and soft materials for energy and biomedical applications. Christopher is a Fellow of the American Physical Society and the North American Thermal Analysis Society (NATAS). He served as the President of NATAS in 2016. Christopher is a recipient of the National Science Foundation Career Award, 3M Non-tenured Faculty Award, and DuPont Young Professor Award among others. He is author/co-author of over 170 refereed journal publications and book chapters.

Contributors

Oscar F. Aguilar Gutierrez Department of Chemical Engineering, McGill University, Montreal, QC, Canada Yong-Ho Ahn Department of Polymer Science and Engineering, Kumoh National Institute of Technology, Gumi, South Korea Kazuo Akagi Research Organization of Science and Technology, Ritsumeikan University, Kusatsu, Japan Shota Ando Department of Applied Chemistry, Meiji University, Kawasaki, Kanagawa, Japan Soma Banerjee Advanced Nanoengineering Materials Laboratory, Materials Science Programme, Indian Institute of Technology Kanpur, Kanpur, India Jin-Hae Chang Department of Polymer Science and Engineering, Kumoh National Institute of Technology, Gumi, South Korea Li Chen The Collaborative Innovation Center for Eco-Friendly and Fire-Safety Polymeric Materials, National Engineering Laboratory of Eco-Friendly Polymeric Materials (Sichuan), State Key Laboratory of Polymer Materials Engineering, College of Chemistry, Chengdu, China Xiaofang Chen Department of Materials Science and Engineering, Drexel University, Philadelphia, PA, USA Suzhou Key Laboratory of Macromolecular Design and Precision Synthesis, Jiangsu Key Laboratory of Advanced Functional Polymer Design and Application, College of Chemistry, Chemical Engineering and Materials Science, Soochow University, Suzhou, People’s Republic of China Donghwan Cho Department of Polymer Science and Engineering, Kumoh National Institute of Technology, Gumi, South Korea Prashant Deshmukh Polymer Program, Institute of Material Science and Department of Chemistry, University of Connecticut, Storrs, CT, USA xiii

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Contributors

Wei Gu Department of Materials Science and State Key Laboratory of Molecular Engineering of Polymers, Fudan University, Shanghai, People’s Republic of China Tae Young Ha Department of Polymer Science and Engineering, Kumoh National Institute of Technology, Gumi, South Korea Edtson E. Herrera-Valencia Laboratorio de Reología y Fenómenos de Transporte, L-7/Primer-Piso; Unidad Multidisciplinaria de Investigación Experimental Zaragoza (UMIEZ), Carrera de Ingeniería Química, Facultad de Estudios Superiores Zaragoza, Universidad Nacional Autónoma de México, Iztapalapa, Mexico Kwang-Un Jeong Polymer Materials Fusion Research Center and Department of Polymer-Nano Science and Technology, Chonbuk National University, Jeonju, Jeonbuk, South Korea Shi Jin Department of Chemistry, College of Staten Island, The City University of New York, Staten Island, NY, USA Ph.D. Program in Chemistry, The Graduate Center of the City University of New York, New York, NY, USA Kamal K. Kar Advanced Nanoengineering Materials Laboratory, Materials Science Programme, Indian Institute of Technology Kanpur, Kanpur, India Advanced Nanoengineering Materials Laboratory, Department of Mechanical Engineering, Indian Institute of Technology Kanpur, Kanpur, India Rajeswari Kasi Polymer Program, Institute of Material Science and Department of Chemistry, University of Connecticut, Storrs, CT, USA Dae-Yoon Kim Polymer Materials Fusion Research Center and Department of Polymer-Nano Science and Technology, Chonbuk National University, Jeonju, Jeonbuk, South Korea Namil Kim Smart Materials R&D Center, Korea Automotive Technology Institute, Cheonan, Chungnam, South Korea Matti Knaapila Department of Physics, Technical University of Denmark, Kgs. Lyngby, Denmark Sandeep Kumar Raman Research Institute, Bangalore, India Department of Chemistry, Nitte Meenakshi Institute of Technology (NMIT), Bangalore, India Christopher Y. Li Department of Materials Science and Engineering, Drexel University, Philadelphia, PA, USA Xiao Li Department of Materials Science and Engineering, College of Engineering and Key Laboratory of Polymer Chemistry and Physics of Ministry of Education, Peking University, Beijing, China

Contributors

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Lalit Mahajan Polymer Program, Institute of Material Science, University of Connecticut, Storrs, CT, USA Hironaga Matsumoto Department of Electronics and Bioinformatics, Meiji University, Kawasaki, Kanagawa, Japan Setsuko Matsumoto Department of Physics, Meiji University, Kawasaki, Kanagawa, Japan Andrew P. Monkman Department of Physics, University of Durham, Durham, UK Kazukiyo Nagai Department of Applied Chemistry, Meiji University, Kawasaki, Kanagawa, Japan Dennis Ndaya Polymer Program, Institute of Material Science and Department of Chemistry, University of Connecticut, Storrs, CT, USA Laurence Noirez Laboratoire Léon Brillouin (CEA-CNRS), Université ParisSaclay, CEA-Saclay, Gif-sur-Yvette Cédex, France Santanu Kumar Pal Department of Chemical Sciences, Indian Institute of Science Education and Research (IISER) Mohali, Mohali, India Qiwei Pan Department of Materials Science and Engineering, Drexel University, Philadelphia, PA, USA Department of Materials Science and Engineering, South China University of Technology, Guangzhou, China Lang Qin Department of Materials Science and State Key Laboratory of Molecular Engineering of Polymers, Fudan University, Shanghai, People’s Republic of China Alejandro D. Rey Department of Chemical Engineering, McGill University, Montreal, QC, Canada Carlos Sánchez-Somolinos Instituto de Ciencia de Materiales de Aragón (ICMA), CSIC-Universidad de Zaragoza, Zaragoza, Spain CIBER in Bioengineering, Biomaterials and Nanomedicine (CIBER-BBN), Madrid, Spain Asit Baran Samui Institute of Chemical Technology, Mumbai, India Shuichi Sato Department of Electrical and Electronic Engineering, Tokyo Denki University, Adachi-ku, Tokyo, Japan Bo-Soo Seo Department of Polymer Science and Engineering, Kumoh National Institute of Technology, Gumi, South Korea Shilpa Setia Department of Chemical Sciences, Indian Institute of Science Education and Research (IISER) Mohali, Mohali, India Zhihao Shen Beijing National Laboratory for Molecular Sciences, Key Laboratory of Polymer Chemistry and Physics of Ministry of Education, Center for Soft Matter

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Contributors

Science and Engineering, College of Chemistry and Molecular Engineering, Peking University, Beijing, China Roman Stepanyan Materials Science Centre, DSM Research, Geleen, The Netherlands Kishore K. Tenneti Department of Materials Science and Engineering, Drexel University, Philadelphia, PA, USA Yingfeng Tu College of Chemistry, Chemical Engineering and Materials Science, Soochow University, Suzhou, China Yu-Zhong Wang The Collaborative Innovation Center for Eco-Friendly and FireSafety Polymeric Materials, National Engineering Laboratory of Eco-Friendly Polymeric Materials (Sichuan), State Key Laboratory of Polymer Materials Engineering, College of Chemistry, Chengdu, China Xiaoming Yang College of Chemistry, Chemical Engineering and Materials Science, Soochow University, Suzhou, China Haifeng Yu Department of Materials Science and Engineering, College of Engineering and Key Laboratory of Polymer Chemistry and Physics of Ministry of Education, Peking University, Beijing, China Yanlei Yu Department of Materials Science and State Key Laboratory of Molecular Engineering of Polymers, Fudan University, Shanghai, People’s Republic of China Lei Zhu Department of Macromolecular Science and Engineering, Case Western Reserve University, Cleveland, OH, USA Tiantian Zhu College of Chemistry, Chemical Engineering and Materials Science, Soochow University, Suzhou, China

1

Introduction to Liquid Crystalline Polymers Soma Banerjee and Kamal K. Kar

Contents Definition . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . An Introduction to Liquid Crystal Concept and Its Origin . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . Liquid Crystalline Polymers: Specialties . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . Classification of Liquid Crystalline Polymers . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . According to Structure . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . According to Phase . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . According to Nature of Mesogens . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . Synthesis of Liquid Crystalline Polymers . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . Physics Behind Liquid Crystalline Polymer Structure . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . Properties of Liquid Crystalline Polymers . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . Processing of Liquid Crystalline Polymers . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . Special Class of Liquid Crystalline Polymers: Liquid Crystalline Elastomers . . . . . . . . . . . . . . . . . . Applications of Liquid Crystalline Polymers . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . Concluding Remarks . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . References . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . .

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S. Banerjee Advanced Nanoengineering Materials Laboratory, Materials Science Programme, Indian Institute of Technology Kanpur, Kanpur, India e-mail: [email protected] K. K. Kar (*) Advanced Nanoengineering Materials Laboratory, Materials Science Programme, Indian Institute of Technology Kanpur, Kanpur, India Advanced Nanoengineering Materials Laboratory, Department of Mechanical Engineering, Indian Institute of Technology Kanpur, Kanpur, India e-mail: [email protected] © Springer Nature Switzerland AG 2020 L. Zhu, C. Y. Li (eds.), Liquid Crystalline Polymers, Polymers and Polymeric Composites: A Reference Series, https://doi.org/10.1007/978-3-030-43350-5_49

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Abstract

Liquid crystalline polymers (LCPs) remain the fascinating class of polymeric material due to the useful combination of physical properties. As the name suggests, LCPs, the class of the macromolecules, possess both the properties of solid and liquid that generate astonishing property in this new and interesting class of materials. These polymers can be of many types depending on the position and type of the mesogenic units in the molecular architecture. LCPs are mainly classified as main-chain, side-chain, crosslinked, etc. made of nematic, smectic, cholesteric, and other mesophases. The orientational properties of LCPs are an important aspect to determine the utility of this particular class of material. The molecules in LCP arrange and align themselves in the longitudinal direction more or less in the transverse direction. This fundamental characteristic of LCPs decides many important properties such as mechanical strength, thermal properties, etc. As a result, these LCPs find applications in several areas such as electrical or electronics, information technologies, medical, aircraft, fiber optics, chemical and domestic equipment, etc., due to the excellent thermal conductivity, good dielectric strength, resistance to solvents, and high dimensional stability. This chapter provides a concise yet informative overview of LCPs starting from its origin, types, synthesis methodologies, essential properties, and application areas. This article also provides a comprehensive overview of the underlined physics behind this structural arrangement leading to molecular anisotropy in the material. The commercial aspects such as processing of this material, how this material differs from the conventional polymer in view of engineering aspects, which make them distinct have been highlighted. Basically this chapter provides a fundamental understanding of this wonderful material in a nutshell. Keywords

Alignment · Anisotropy · Birefringence · Chemical resistance · Cholesteric · Dielectric constant · Director · Liquid crystals · Liquid crystalline polymers · Lyotropic · Mesogens · Mechanical strength · Molecular domains · Modulus · Mesophase · Monodomain · Nematic · Network structure · Orientation · Processing · Smectic · Spinning · Supramolecular structure · Thermal expansion · Thermotropic

Definition Liquid crystalline polymers (LCPs) with specific mesogen features are the polymers of specific class having distinct properties compared to the conventional polymers. They exhibit a combination of properties of liquid and solid crystal leading to distinct behavior. These polymers find applications in numerous areas of research from medical, chemical, electrical, optical, mechanical, and many more due to the tunable orientational properties that distinguish them from the traditional class of polymers.

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Introduction to Liquid Crystalline Polymers

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An Introduction to Liquid Crystal Concept and Its Origin Academic studies on liquid crystals have started more than 120 years ago, which have now taken the interest of the scientific community for their possible vast practical applications (Reinitzer 1989; Dunmur and Sluckin 2014). As the name suggests, liquid crystal compounds can combine the properties of a solid crystalline material and a liquid. Just like a liquid, they can flow and take the form of the container where it holds; however, at the same time, they exhibit anisotropic behavior similar to solids in an astonishing way. Hence, the physical characteristics of these materials are expected to be different in different directions named as anisotropic liquids sometimes. Thermodynamically, a liquid crystalline state of a material is a stable phase, which is an intermediate of the amorphous and crystalline state of a solid. Therefore, this phase of the materials is often named as the mesomorphic state, also called the mesogens originally adapted from the Greek word “mesos” meaning intermediate (Goodby et al. 2014). Once again the origin of liquid crystals is way back to the year 1888 when Austrian chemist Friedrich Reinitzer has observed a peculiar behavior of a solid after subjecting it to the different temperatures (Reinitzer 1989). He has noticed that at a particular temperature, cholesteryl benzoate solid becomes hazy liquid and again is converted to a clear liquid once the material is heated to a higher temperature. During cooling, he has observed that the liquid again passes through two different colors before, finally, it achieves the white solid nature. This material also exhibits different melting points, which is quite surprising. Later on, Lahmann analyzed the material and found the presence of multiple small crystallites with improper boundaries indicating that the first intermediate fluid remains crystalline by nature and expecting this to be a new state of a matter. After prolonged studies and analysis, Lahmann named his finding as liquid crystals (Lehmann 1900). Later in 1850, Heintz has also observed the fact that natural fats show two different melting points (Ramberg 2013). A representation of molecular ordering has been displayed in Fig. 1. By nature, liquid crystals are composed of small crystalline elements that are suspended in the liquid phase. Unlike the pure materials, they possess two melting points. The individual molecules of the material that are capable to form mesophases

Fig. 1 A representation of molecular order in (a) crystalline, (b) liquid crystalline, and (c) isotropic liquid

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Fig. 2 Orientational order of mesogens in liquid crystal molecules: (a) nematic, (b) smectic A, and (c) smectic C

are named as mesogens. Mesogens are small organic molecules capable to arrange concisely in different types of orders or organizations. However, all these mesogens do not participate in organization or ordering and hence generating an astonishing material having a combination of both solid and a liquid. For example, gelatin shows mesophasic nature. The functional behavior of a liquid crystal lies in the organization or order of the mesogens. Since, liquid crystals remain in two states, in a liquid state, the mesogens become arbitrarily oriented without any directionality and form an isotropic fluid. Again, in the solid state, the same mesogenic molecules are in a highly ordered state with no translational degree of freedom. For liquid crystal materials, the difference lies in the sense that the non-symmetrical mesogen molecules are self-aligned in a particular axis called the director (Fig. 2) (Ermakov et al. 2015). This positional arrangement is the most critical for the formation of the nematic and smectic structures as well as solid phases. A nematic liquid crystal molecule retains their directional orientation; however, it contains some sort of freedom of movement within the liquid crystals as depicted in Fig. 2a (Kumar 2001). In the smectic phase, the liquid crystals are arranged in such a way that their principal axis is oriented parallel to the center of mass of the material in one of the planes, thus exhibiting both positional and directional order in these molecules, as represented in Fig. 2b and c (Ermakov et al. 2015). Hence, all the properties of interest of liquid crystals such as optical, magnetic, and electrical are direction oriented or anisotropic.

Liquid Crystalline Polymers: Specialties Later on, the researchers have observed that liquid crystals may also produce polymeric molecules consisting of repeating monomer units, which are connected to form long-chain molecules. The primary units of the main chain are connected to each other through a flexible linker of varying chain lengths. These polymeric chains

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are capable to assemble to generate LCPs in a similar way liquid crystal does with a single mesogen molecule. They have a fine combination of physicomechanical properties and extraordinary optical properties similar to liquid crystals. The LCPs have great potential and a matter of interest to the modern research field due to the tunable ordered supramolecular structures. The structural organization plays a major role in deciding the physicomechanical properties of these polymers since these can be easily altered by external mechanical and electromagnetic fields, thermal effects, light irradiations, etc. Figure 3 shows the general structures of main- and side-chain LCP structures with flexible and semi-flexible crosslinkers. Like general polymers, LCPs can also be of two chain types, rigid and flexible. The structural design in LCPs can be obtained by two routes (Shibayev and Byelyayev 1990). In one case, rigid chain macromolecules produce liquid crystal phases by the spontaneous orientational ordering of the long-chain molecules, or in another case a flexible chain macromolecule can be used, which can be modified further by inclusion of low molecular weight mesogens or rod-like mesogens in the main chain of the polymer as shown in Fig. 4 (Moeller and Matyjaszewski 2012). Rigid LCPs appeared to be less important due to the high melting points, and the melting of these polymers may occur at a high temperature, which is close to the thermal degradation stage of the polymer (Donald et al. 2006). However, dissolution in the proper solvent may produce stable lyotropic liquid crystals leading to the formation of high strength synthetic fibers (Khokhlov and Semenov 1981). In the second case, thermotropic LCPs come into the picture that is considered to be more attractive. For these LCPs, chemical bond formation takes place between flexible and rigid or mesogenic segments in a macromolecule. Again, this may lead to the formation of two types of LCPs, one with linear structure and other with a branched structure. In case of linear LCPs, the mesogenic groups have been added in the main polymer chain, whereas, in the branched LCPs, the mesogenic units are incorporated in the main polymer chains by chemical bonding with a flexible spacer, Fig. 3 Examples of mainand side-chain LCP structures with flexible, semi-flexible crosslinkers. (Reprinted with permission from Cresta et al. 2018)

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Fig. 4 Macromolecules of different chain rigidity: (a) flexible polymer chain, (b) rigid polymer chain, (c) LCPs containing mesogens in the main chain, and (d) LCPs containing mesogens in the side-chain (comb-shaped LCPs)

generally alkoxyl or alkyl types leading to the formation of branched LCPs commonly known as comb-shaped LCPs. LCPs with mesogens in the main chain already find applications in numerous fields, development of strong fibers based on lyotropic liquid crystalline polyamides, such as Kevlar, Twaron, etc. The strength and modulus of these LCP-based fibers are about 2–2.5 and 10–20 times better as compared to those made of yarns of aliphatic polyamides and about 2–4 times greater compared to steel and glass fibers (Shibaev and Bobrovsky 2017). This sudden improvement in mechanical properties is due to the packing of mesogenic units in the liquid crystal state leading to the generation of systematic regions with parallelly oriented polymer chains during yarn formation or spinning of the fiber. Among the liquid crystal-based plastics, Xydar and Vectra are of great technological importance due to the extremely beneficial mechanical properties used as self-reinforced plastic materials (Kar and Otaigbe 2004). The structural arrangement of the mesogenic units takes place at the time of melt spinning during extrusion of the melt through the aperture of the spinneret. New-generation plastics based on liquid crystalline polyester are also prepared by the melt extrusion process. These superplastics are self-reinforcing in nature as described earlier and are of high modulus and strength of about 60–70 GPa and 700 MPa, respectively, with a very low elongation at break of 1.5–2% (Shibaev and Bobrovsky 2017). Liquid crystal-based polyesters are also of low coefficient of thermal expansion (
106 K1) that are comparatively lower than that of non-crystalline of the same type (104 K1) (Shibaev and Bobrovsky 2017). As mentioned earlier, LCPs can be broadly classified as thermotropic LCPs and lyotropic LCPs depending on the presence of liquid crystalline phases either in the melt or in the solution form. Lyotropic LCPs are of good liquid crystalline nature, and the extent of crystallinity is dependent on the nature and temperature of the solvent and concentration of the polymer. These materials are incapable of exhibiting liquid crystallinity in the melt state since they degrade before melting and exhibit transition of phases through addition or removal of solvent molecules (Gray et al. 2009). For thermotropic LCP materials, the transition among the phases is governed by a thermal process and controlled by the thermal history of the material and melt temperature. The thermally triggered liquid crystal mesogens form liquid crystal phases from the crystalline melting point to the isotropic temperature of the material (Gray et al. 2009).

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Fig. 5 Examples of basic shapes forming LC phases (rod, disk, board, and banana shapes), different types of molecular architectures possible in LCP structures

Thermotropic LCPs are anisotropic melts by nature and possess comparatively low viscosity that has grabbed the attention of researchers due to the excellent chemical resistance, outstanding mechanical properties, low permeation of liquid and gases, dimensionally stable with low coefficient of thermal expansion, etc. They can be processed easily by injection molding and extrusion (Kar and Hodzic 2011; Kar 2011). Thermotropic LCPs can again be oriented to form fibrous structures due to the extraordinary strength and stiffness originated from the presence of rigid rod-like molecules ensuing in self-reinforcing nature (He and Bu 1994). Hence, thermotropic LCPs are of great technological importance both in neat form and also in thermoplastic composites in applications especially related to high-performance engineering fibers (Kim and Kim 2006). Figure 5 shows the possible structures of LCPs containing rigid mesogens in the main chain, side chain, and their combinations giving rise to linear or branched LCPs.

Classification of Liquid Crystalline Polymers Based on the discussion on the earlier section, an attempt is made for better understanding to classify LCPs according to structure, mesogenic group, and phases as follows.

According to Structure Amphiphilic LCPs: These LCPs are soluble in solvents or water or both and named as lipophilic and hydrophilic, respectively. These can be further subdivided into the

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class of ionic (cationic or anionic) and non-ionic. The lyotropic LCPs are formed by direct participation of amphiphilic molecules with water leading to the formation of mesophases. Non-amphiphilic LCPs: These can be referred to as the nonpolar or organic polymers of less polarity possessing high anisotropy in its geometry due to the inherent rod or disk-like structures. These LCPs generally form mesogenic phases after melting defined as thermotropic LCPs.

According to Phase LCPs can be classified into two main groups based on the liquid phase as thermotropic and lyotropic. Thermotropic LCPs are generated by the application of heat when structural ordering in pure molecules takes place, whereas in the case of lyotropic LCPs, mesogenic ordering of phases is observed in the presence of suitable solvent molecules. Thermotropic LCPs can again be subcategorized as enantiotropic and mesotrophic phases. In the case of enantiotropic LCPs, the formation of a liquid crystalline phase takes place during both heating and cooling cycles, whereas mesotrophic LCPs are formed from the isotropic liquid and are stable on a supercooled stage. The mesotrophic LCPs can be subdivided into three more groups such as smectic, nematic, and cholesteric as described earlier.

According to Nature of Mesogens Depending on the nature of mesogenic groups, LCP structures may differ to a great extent, and this can be the main-chain LCPs and side-chain LCPs. As the name suggests, the main-chain LCPs are those materials, where the mesogenic units are itself a part of the main chain of the polymer, whereas in case of side-chain LCPs, the mesogen units are attached like a side chain or pendant with a flexible spacer molecule in the main polymer backbone. Structurally, a main-chain LCP is formed when rigid groups are attached to the comparatively flexible polymer chain. These can be formed from a rigid and rod-like monomer molecule. In another case, the main-chain LCPs can be generated due to the direct incorporation of mesogenic groups into the main polymer chain. The mesogens so attached work just like some stiff areas inside the first group. The liquid crystal behavior is observed due to the structural restriction imposed by the mesogenic aromatic rings. The side-chain LCPs are composed of three main structural units, namely, the backbone, spacer, and mesogen unit. The backbone provides the spine, where the side chains are attached. The structure of the main chain also decides an important role in determining the fact that whether the liquid crystalline structure will be formed or not. The most important part has been played by the mesogen units attached to LCPs, since the alignment of structural units generates the liquid crystal behavior. In general, the mesogens are made of rigid core or two or more aromatic groups clubbed with a functional group. Figure 6 represents some possible structural arrangements of the main-chain, side-chain, and crosslinked LCPs formed from mesogenic units and flexible spacers.

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Fig. 6 Schematic of (a) main-chain LCP, (b) side-chain LCP, (c) side-chain mesogens attached to the spacer, (d) side-chain mesogens attached to mesogens, and (e) crosslinked LCPs

Synthesis of Liquid Crystalline Polymers There are some principles for the synthesis of side chain bearing comb-shaped LCPs that can be broadly divided into three classes. Examples of the chemical structure of few monomers used for the synthesis of LCPs are represented in Fig. 7. First method remains the synthesis of monomers bearing mesogenic groups followed by homo- or copolymerization with other mesogenic or non-mesogenic units that finally leads to the synthesis of comb-shaped LCPs. In the second method, LCPs can be synthesized chemically binding the mesogen units to the already prepared polymer main chain leading to the creation of side-chain LCPs. However, this method is less preferred compared to the first one, since this method may generate compositional inhomogeneity in the final LCPs and, in addition to that, other constraint remains a difficulty in obtaining the polymers with a complete degree of substitution. Liquid crystal polysiloxanes are prepared mostly by this method, where hydrosilylation reaction between polysiloxanes and mesogen-type monomers is carried out as represented in Fig. 8. where R is the mesogen. The ease of this kind of reaction leads to the formation of LCPs with a certain degree of polymerization leading to the commercialization of polysiloxanes. This type of synthesis methodology is also utilized for the synthesis of liquid crystal-based elastomers and dendrimers based on the chemistry of carbosilane derivatives. In the third synthesis method, LCPs are synthesized by using condensation polymerization. Polycondensation reactions can generate both main- and

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HO

O

O

C

C

HO

O C

OH

OH

HO

Hydroquinone

Terephthalic acid OH C

O C

O OH

HO

HO

C O

Isophthalic acid

Naphthalene-2,6-dicarboxylic acid

OH

4-(4-hydroxyphenyl)phenol

Fig. 7 Chemical structures of monomers producing LCP structures

Fig. 8 Reaction between polysiloxanes and mesogen

side-chain LCPs. These three general methods lead to the formation of covalently bonded liquid crystal homo- and copolymers. Liquid crystalline polymer networks (LCNs) are the recent class of LCPs made of loosely crosslinked LCPs with mesogens in the polymer structures either in its main or side chain. When mesogens are aligned in a uniaxial direction, the LCN can reversibly transit between shrinkage in isotropic state and expansion in liquid crystal state. Hence, these materials will behave differently in an aligned direction and perpendicular direction leading to a rapid and extreme shape change. This property of LCN makes them useful in several applications such as robotics, sensors, actuators, optics, biomedical applications, etc. (Ohm et al. 2010; Kar 2016; Ditter et al. 2017). Other important areas of application of LCN belongs to energy generators, soft robotics, motors, actuators, etc. (Tang et al. 2015; Wie et al. 2016). LCN-based actuators are prepared by the alignment of mesogens via application of mechanical force or electric or magnetic field. Afterward, the liquid crystal alignment gets fixed by the crosslinking of the polymer chains (Yoon et al. 2018). LCN actuators can develop an intricate and reversible change in shape due to the phase transitions of the mesogen units achieved via controlling the alignment of liquid crystals through the disbursement of the crosslinkable area or via formation of patterned stimulations

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(Saed et al. 2019). A common way to prepare well-organized LCN is via photoinitiated polymerization of liquid crystal monomers (Broer et al. 1989; Broer 1993). The monomers used for the development of LCPs by photoinitiated polymerization have a structure similar to conventional liquid crystals. They possess anisotropic shape with a tough central core of rod or disk-like nature. In addition to generic structural arrangement, the liquid crystal monomers possess groups that can be polymerized. The polymerizable groups can be positioned either on the lateral side or at far ends. The liquid crystal monomers may possess one, two, or more functional groups. A representation of common liquid crystal monomers and liquid crystal networks has been displayed in Fig. 9. Photopolymerization is another most commonly used process for acrylates and methacrylates via exposure under UV radiation at a high rate of photopolymerization (Fig. 9). A small concentration of photoinitiator is needed for the progress of the reaction. The use of small pendant groups is of most interest due to the provisions of low melting temperature leading to easy processing and reduced thermal polymerization. An increment in spacer length leads to the formation of a stable smectic phase. During photopolymerization, all the components are dissolved in strong solvents like xylene, THF, dichloromethane, etc., which are evaporated at the later stage of polymerization. Thin films are formed either by processing from solution or melt stage. In general, pre-treated substrates such as glass coated with thin polyimide coating are used to generate molecular orientation of liquid crystal monomers

Fig. 9 Examples of common liquid crystal monomers and liquid crystal networks. (Reprinted with permission from Liu and Broer 2014)

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perpendicular to the substrate material. Complex structures can be formed when liquid crystal melts are passed through the capillary filling. This leads to the formation of splayed and twisted nematics. Figure 10 shows the scanning electron microscope image of a fractured film of a chiral nematic liquid crystal acrylate polymeric network showing the molecular orientation present therein. When the desired structure is generated, the liquid crystal monomers are exposed to light to initiate the polymerization process. Depending on the types of photoinitiators, the polymerization may proceed under UV or visible light source. For liquid crystal acrylate polymerization, the free radical polymerization proceeds by avoiding oxygen attack. Photopolymerization is, in general, a high-conversion method, specifically when performed at the elevated temperatures. However, if the curing process is pursued at room temperature, vitrification may proceed during polymerization. This restricts the mobility of the monomers and further reaction is hampered. In these cases, post-curing at or around glass transition temperature may be helpful. In practice, curing at 120  C for 10 min after the UV exposure is sufficient for the full polymerization process (Liu and Broer 2014). Figure 11 represents a typical example of surface dynamics of patterned nematic networks. Mostly all commercial LCPs contain p-hydroxybenzoic acid ( p-HBA) as a monomeric unit in its structure. These LCPs are synthesized by condensation of p-HBA with other monomers. A homopolymer composed of only p-HBA generates liquid crystals of flowability below 500  C (Linstid et al. 2000). The processing temperature can be reduced further by the addition of different monomers such as bisphenols, terephthalic acid, isophthalic acid, hydroquinone, etc. As, for example, a

O

CH3

H

O

O

O O

O

O O

O H CH 3

O

p = 180 nm

Fig. 10 Scanning electron microscope image of a fractured film, cross section of a chiral nematic liquid crystal acrylate polymeric network. (Reprinted with permission from Liu and Broer 2014)

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Fig. 11 Surface dynamics of patterned nematic networks (a) representation of ordering and characterization by polarization microscopy, (b) surface profile as prepared and during exposure to UV light, (c) mechanism of surface deformation on UV exposure, and (d) three-dimensional images of surface topographies of original, UV irradiated, and after removal of UV source. (Redrawn and reprinted with permission from Liu and Broer 2014)

high-melting polyester of poly(4-hydroxybenzoic acid) has been formed by condensation of p-HBA monomers, while copolymers are formed by the reaction of p-HBA with terephthalic acid and bisphenol A. Figure 12 represents the chemical structures of some industrial-grade LCPs.

Physics Behind Liquid Crystalline Polymer Structure The orientational properties of LCPs are an important aspect to determine the utility of this particular class of material. The molecules in LCP arrange and align themselves in the longitudinal direction more or less in a transverse direction. This fundamental property of LCPs decides many important properties such as mechanical strength (Hamley et al. 1996). The degree of orientation is usually described by a unit vector, “n” (Allen et al. 1996). This vector “n” is commonly defined as the director or optical axis (Marrucci 1996). The director “n” is determined by the action of the flow of the molecules or some weak forces such as magnetic and electric forces, and it is not related to the individual orientation of the molecules; rather it is dependent on the average molecular orientation inside the LCP. This vector is further quantified and characterized by scalar quantity “S” also known as the order parameter or Herman’s orientation function and calculated using Eq. 1 (Pavel et al. 2005):

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S. Banerjee and K. K. Kar O C

O

C

O

Vectra LCP

O O C

O

O

O

O

O

C

C

Xydar LCP O O

C

O

H2 C

H2 C

O

O

O

C

C

Rodrun LCP Fig. 12 Chemical structures of industrial-grade LCPs

S ¼ 1=2 3 < cos 2 θ > 1




ð1Þ

where θ represents the angle of orientation between the mesogen of the polymer chain and the director “n.” The bracket around the cos2θ indicates the thermal average over all the unit molecules. When the value of “S” becomes zero, it indicates total isotropy in the material, whereas when S becomes 1, it indicates perfect ordering among the molecules, where all the molecules are parallel to each other (Ten Bosch et al. 1983). This ideal ordering of the molecules is only possible near absolute zero temperature if the material does not freeze. Hence, the order parameter of the material is inversely proportional to the kinetic motion of the molecules. The exact value of S remains a compromise between the ordering effect of the mesogens and the contribution of temperature to the disordering as well (Klein et al. 1996). In general, S decreases with an increase in temperature and varies from 0.43 at Tc, clearing temperature, to about 0.8–0.9 at lower temperatures (Nyden and Gilman 1998). For LCP having a combination of main- and side-chain macromolecules, the order parameter is a combined contribution from both main and side chains. The order parameter is also known as the anisotropy factor that determines the properties of the material depending on its direction of measurement. Liquid crystals and LCPs both exhibit different mesophases such as nematic, smectic, columnar, cholesteric, etc., due to different degrees of molecular ordering in

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the material. In general, rod-like LCPs form nematic, cholesteric, and smectic phases, while disk-like molecules show columnar and nematic mesophases. Nematic remains the most common and least organized liquid crystalline phase. This phase results when the molecules are ordered in one dimension only, and the ordered phases remain parallel to each other. Smectic is another well-known liquid crystalline phase formed when the molecules are arranged in the long-range. In this particular case, the molecules remain parallel and ordered in layered arrangement one over the other in the form of stacks. When an LCP contains both phases, the smectic phase arises at a lower temperature. In the other phase, the columnar phase is made during the heating or cooling process of the molecules that are disk-shaped in nature and capable to be packed together in the form of cylinders or columns similar to the way to stacking of coins. The individual molecules within these layers of stacks may be ordered or disordered. The columns can again be combined together in the form of hexagonal or orthogonal lattice forms (De Gennes and Prost 1993). LCPs may form twisted or chiral nematic phases named as cholesteric liquid crystalline phase. Cholesteric phases are formed when nematic phases orient themselves in the form of layers and each layer is again twisted with respect to one and another. The director “n” is considered to be variable in space and twisted periodically in an axis normal to the vector “n” leading to the formation of helical form. The pitch of the helical form is defined as the distance, over which the vector goes 360 . This particular arrangement of phases may result in unique optical properties. This cholesteric phase owns orientational order; however, long-range or positional orders remain absent in these molecules. Figure 13 shows the position of the director and the changes in the phases during heating of a thermotropic LCP. Nematic, cholesteric, and smectic are the phases that are commonly observed in LCPs, whereas other high-order phases are less commonly observed in LCPs.

Fig. 13 Pictorial representation of the effect of heating in thermotropic LCP and change in phases; ϑ represents the angle between the long axis of director and individual molecules; n stands for director. (Redrawn and reprinted with permission from Dierking and Al-Zangana 2017)

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Properties of Liquid Crystalline Polymers LCPs show excellent mechanical and physicomechanical properties due to the combination of liquid and crystalline nature. They have potential in photoresponse-related applications, haptic displays, flow control, catalysis, optics/ photonics, etc. LCNs exhibit negative thermal expansion and multiple phase transformation in the same material. The photoresponsive behavior in LCNs is due to the anisotropic organization of mesogenic moieties. Azobenzene-containing LCNs are recently utilized as shape memory polymer and adaptive material (Harris et al. 2005; Lee et al. 2011, 2012). McConney et al. have reported a photoresponsive LCN, where all the surface features have been initiated by the application of photons (McConney et al. 2013). One of the striking features of liquid crystal materials is that they have the capability to arrange direct profiles into intricate patterns. The direction-oriented patterning is possible by the use of photo-aligned surfaces based on azobenzene material. In an interesting study, Broer and group have fabricated freestanding films composed of three-dimensional molecular ordering via photoalignment of polymerizable liquid crystals (de Haan et al. 2012). The film so obtained upon heat treatment deforms to cone and saddle forms. Similar to low molecular weight liquid crystals, the liquid crystal monomers are quite birefringent and have a high refractive index when the measurement is carried out under light polarization in a direction parallel to the director. When the measurement is conducted in a direction orthogonal to the director, the refractive index becomes much lower. The monomer behaves similarly as that of low molecular weight nematic liquid crystals and possesses a huge dependence on temperature when transition proceeds from nematic to isotropic phases. Upon heating of the polymer to a temperature near thermal degradation, no more new isotropic phase is formed, and the birefringence is only nominally affected. This remains a common behavior for all the LCPs based on diacrylates. The temperature dependence on the optical properties of the LCPs is much dependent on the spacer chain length and crosslinking density of the polymer. LCPs of aramid types exhibit good elastic modulus and tensile strength as compared to that of networks formed from LCPs of the same kind. The liquid crystal networks exhibit modulus in the range of GPa, and the tensile strength remains in the range of 10–100 MPa (Liu and Broer 2014). The modulus remains anisotropic in nature; however, it is only three times higher when the measurement is conducted in a direction along the director as compared to that perpendicular to it. Anisotropic thermal expansion of LCN remains much interesting compared to its strength or modulus parameters (Broer and Mol 1991). The majority of the covalent bonds remain in the direction of the director, which leads to lower linear thermal expansion in the direction parallel to the director. Again, during thermal treatment, as the material passes above the glass transition temperature, the order parameter is decreased further leading to a contraction in the material upon further increase in the temperature. The linear coefficient of thermal expansion remains negative in the direction parallel to the orientation direction. The thermal expansion becomes much large when the measurement is conducted in the perpendicular direction and shows

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an increasing trend with temperature. Above the glass transition temperature, the anisotropic property of thermal expansion becomes larger. The influence of temperature of polymerization and the presence of substituents play a minor role in this respect. A high polymerization temperature in general leads to the formation of less ordered network structure and, hence, the anisotropy in the material decreases. In general, LCPs are of high mechanical strength at high temperatures, superior chemical resistance, good weatherability, and flame-retardant properties. They may create several forms via sintering and molding. LCPs possess a high coefficient of thermal expansion along the Z-axis. LCPs are of exceptional chemical resistance at elevated temperature and in presence of strong acids, bases, and aromatic and halogenated hydrocarbons. The dimensional stability of LCP is quite good even in boiling water. Polar and bowlic LCPs are by nature ferroelectrics with a very low reaction time compared to conventional liquid crystals making them a material of choice for ultrafast switches. Bowlic polymers are columnar in shape and can be used to form ultrahigh superconductors. Table 1 shows some important properties of few industrial-grade LCPs. LCPs are known for high modulus thermoplastic matrix. It is also reported that the addition of LCPs to a traditional thermoplastic matrix reduces the viscosity of the compounding process, and hence the processability could be improved to certain extent. The LCP/thermoplastic blends also have other advantages over the standard glass-fiber-reinforced composites in addition to reduced melt viscosity and slashed energy consumption in processing. Few important LCP-based composites along with properties are mentioned in Table 2.

Processing of Liquid Crystalline Polymers LCPs are generally processed and molded by five ways: melt spinning, pressing, injection molding, extrusion, and coating (Kar and Hodzic 2011; Kar 2011). Aromatic LCPs composed of copolyesters exhibit exceptional orientational behavior and properties in its solid state. In an interesting study, the orientation of molecules is examined under shear and elongation flow/deformation. The study exhibits that high molecular orientation can be observed under elongation, but no such observation can be seen in case of shear force. Rod-shaped molecules from isotropic solutions are oriented in the flow direction, whereas, in the case of shear stress, the molecules will revolve intermittently. Again, if the molecules are under an anisotropic state in the melt or solution phase, where the molecular domains are already oriented, the application of shear stress will not affect the orientation of stabilized molecules. However, in the case of elongational flow, the molecular domain will preferably orient and stretched in the direction of flow yielding monodomain state formation (Ide and Ophir 1983). During the injection molding of the LCPs, a distinct core and skin morphology has been reported. This can be explained by the quenching process at the walls and elongational flow history of the material. The molten polymer when being injected into the cavity touches the flow front and experiences a strong elongational flow and

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Table 1 Properties of industrial-grade LCPs Trade name RODRAN Aromatic copolyester containing 60% p-hydroxybenzoate, hydroxybenzoic acid, and 40% poly(ethylene terephthalate)a VECTRA Aromatic copolyester containing 73 mol% hydroxybenzoic acid and 27 mol% hydroxynaphthoic acidb

Manufacturer Unitika

Grades LC 3000

Specifications TS: 120 MPa HDT 150  C), where the entropy has a greater influence, the whole polymer chains act as cylinders and organize in a columnar nematic phase, probably behaving as self-compacting chains. This is supported by the appearance of the isotropic phase when cooling the columnar nematic phase to around 150  C. Further cooling the isotropic phase to ~130  C, a Colh phase appears, driven by the formation of TP π-stacks (Zhu et al. 2014). Ring-opening metathesis polymerization has been employed to prepare monodispersed TP-functionalized poly(norbornene)s (PTPNB-m) and poly(butadiene)s (PTPBD-m) with the structures shown in Chart 7. It has been shown that both PTPNB-10 and PTPBD-10 form a disordered Colh phase without obvious π-stacking order and very low clearing point, while PTPNB-5 and PTPBD-5 are amorphous. Hydrogenation of the polybutadiene main-chain (leads to the formation of PTPE-m) does not significantly affect the stability of the mesophase (Weck et al. 1997). PTPA-m (Yu et al. 2013) and PTPOA-m (Xing et al. 2008) are TP-functionalized polyacetylenes. PTPA-1 and PTPA-3, with short spacers, form a Colh phase that consists of cylindrical chains having TP units jacketing around the polyacetylene main-chain. In contrast, with a much longer spacer, the TP π-stacking dominated the self-assembly of PTPOA-m in their Colh phase. Column-forming TP side-units and layer-forming poly(3-alkylthiophenes) are combined together in (PT5T)n, (PT10T)n, and (PT15T)n (Tahar-Djebbar et al. 2011; Zeng et al. 2014). (PT10T)n and (PT15T)n show a rectangular lamello-columnar mesophase. The lamellar order arises from the phase separation of rigid polythiophene backbones from the matrix and columnar order arising from π-stacking of TP units, with TP columns parallel to the polymer backbone direction. Similarly, (PT5T)n can also organize into a mesophase with a lamello-columnar order. However, the 2D order in the phase is of short range; therefore, the phase can be assigned as a lamello-columnar nematic phase.

Columnar LC Phase-Forming Polymers with Pendent Tapered Minidendrons Although tapered minidendrons are not discotics, they can be considered as columnar mesogens as many of them are capable of forming a columnar mesophase (Rosen et al. 2009). A Colh phase has been observed in PEI-m, a series of poly/oligoethyleniminesbearing tapered minidendron side-groups (Fischer et al. 1995; Seitz et al. 1996; Percec et al. 2001). The general structure of PEI-m and MW-dependent clearing point of the Colh phase of PEI-8 are shown in Fig. 8. Such a weak MW dependence of the clearing point, along with the practically MW independent Colh lattice

136 120

Clearing Point (°C)

Fig. 8 Structure of PEI-m and DP-dependent clearing point of PEI-8

S. Jin

100 80 60 40 20 0

PEI-m, m = 8 - 13 0

50

100

150

200

Degree of Polymerization

parameter, suggests that the phase formation is dominated by the minidendrons. This is also supported by the fact that the change of the type of the minidendrons strongly affects the stability of the Colh phase. The use of one-tailed or three-tailed minidendrons or even 3,5 two-tailed minidendrons destroys the ability of the polymer to show a mesophase. It was suggested that most likely each polymer chain forms a cylinder in the Colh phase, with tapered minidendrons radically surrounding the helical main-chain. PECH, poly(epichlorohydrin), randomly functionalized with a three-tailed tapered minidendron (Chart 8), can exhibit a columnar LC phase when the degree of substitution is 48% or higher (Ronda et al. 2003; Giamberini et al. 2005). The isotropization temperature increases with the degree of substitution. X-ray diffraction results suggest that the benzene rings are tilted at 20–30 with respect to the column axis in the mesophase. The same minidendron was also attached to the poly (7-oxanorbornene) and polymaleimide main-chains (Percec and Schlueter 1997). Both polymers exhibit a Colh phase, with the polymaleimide main-chain taking 7/2 helical conformation while poly(7-oxanorbornene) main-chain taking 3/1 helical conformation. A number of minidendrons with added rigid-rod units have been reported to induce a columnar LC phase in polymers (Rosen et al. 2009). Chart 9 depicts some of such polymers. While both 3,4-PS and 3,4,5-PS form a Colh phase, 3,5-PS is an amorphous liquid, suggesting that the columnar mesophase is indeed dominated by the selfassembly of minidendrons (Percec et al. 1998b). The impact of the length and the nature of the spacer connecting the minidendron substituent and the main-chain on the columnar mesophase was investigated in BPMA-m, m = 1–4 (Percec et al. 1993). It was observed that all polymers form a Colh phase, and the lattice of a minidendron-jacketed polymethacrylate is slightly larger than that of the corresponding minidendron. An increase of spacer length in a polymer increases the diameter of columns and reduces the clearing point of the Colh phase. In addition, the incorporation of a larger periphery aryl unit such as naphthalene or biphenyl into BPMA-4 leads to a similar Colh phase with a slightly larger column (Percec et al. 1998c). More detailed X-ray analysis suggested that in the Colh phase of BPMA-4, minidendrons organize around the polymer main-chain in an 81 helical order (Rosen et al. 2009).

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PECH

137

POBN

PMAI

Chart 8 Structures of columnar LC phase-forming polymers carrying a three-tailed minidendron

m = 1: R1 = -C12H25 m = 2: R2 = -C5H11 m = 3: R3 = -CH2CH(CH3)C2H5 Y =CH2, C=O

m = 1: R1 = -C12H25 m = 2: R2 = -C5H11 m = 3: R3 = -CH2CH(CH3)C2H5

m = 1: A1 =

L1 = C3H6

m = 2: A2 =

L2 = C3H6

m = 3: A3 =

L3 = C11H22

3,4-PS 3,4,5-PS

BBPBPMA-11

3,5-PS

BPMA-m, m = 1-4

BPPMA-4

PVE-1

PVE-2

NPMA-4

Chart 9 Structures of columnar LC phase-forming polymers carrying minidendrons with added rigid-rod units

MW dependence of dendron-jacketed polymers has been explored using polyvinyl ethers with narrow MW distribution. PVE-1 always forms a Colh phase, regardless of the DP (3.7–16). It was suggested that with a DP < = 4, the dendrons self-assemble into discs that stack to form a Colh phase. With a DP > 4, each polymer chain forms a column, with the main-chain threads through the center (Percec et al. 1992a). In contrast, PVE-2 forms a Colh phase above the melting point of the underlying crystalline phase at low DP (3.2–3.9), where columns form in the same way as in low DP PVE-1. However, PVE-2 is amorphous at DP > 5. At DP 5, the polymer displays a Colh phase and a reentry isotropic phase (Percec et al. 1992b).

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Chart 10 Structures of columnar LC phase-forming polymers carrying dibenzo[a,c]phenazine (left) and perylene diimide mesogens (right)

Other Discotic Side-Chain Polymers Tetrahexyloxydibenzo[a,c]phenazine-based discotic mesogen was randomly attached to the backbone of polyacrylamide (Chart 10). With 50% mol repeating units modified, the polymer exhibits a Colh phase, as suggested by X-ray diffraction and texture results. In each column, the inter-core spacing is about 0.34 nm (Lee and Huang 2012). Perylene diimide functionalized polyacrylate demonstrated an oblique 2D columnar LC phase with two mesogens in every unit cell. Along the column direction, the perylene cores packed in a disordered fashion with an average intercore distance about 0.34 nm. In comparison, the small molecular model compound does not exhibit a LC phase. It is believed that the attachment of the intrinsic disordered atactic polyacrylate chains suppressed the crystallization of perylene units and leads to the formation of columnar LC phase (Kohn et al. 2012). Although monomeric hexabenzo[bc,ef,hi,kl,no,qr]coronene(HBC) derivatives have shown a strong tendency to form a columnar phase, columnar LC phaseforming covalent side-chain HBC polymers have not been reported, to the best of the author’s knowledge. On the other hand, columnar LC phases have been observed in ionic complexes of anionic HBC moieties and cationic polymers. The complex between HBCA1, a carboxyl-functionalized HBC, and a hydrophobic modified polyethyleneimine, as shown in Chart 11 (left), exhibits a rectangular columnar LC phase below 70  C and a Colh phase above 70  C. Note that the free HBC itself exhibits a low-temperature rectangular phase and a high-temperature Colh phase. One interesting observation is that in the complex, the intracolumn order is better than the free HBC (Thünemann et al. 1999). Chart 11 (middle) also depicts a polysiloxane/HBCA1 complex that exhibits two oblique columnar LC phases. In the lower-temperature phase, the HBC cores are tilted with respect to the column axis. In the higher-temperature phase, the HBC plane is perpendicular to the column axis. An unusually long correlation of HBC units was observed (Thünemann et al. 2000b). Columnar mesophases have also been observed in the complex of poly[ethylene oxide]-block-poly[L-lysine] (PEO-PLL) and HBCA2 (Chart 11, right), a carboxylfunctionalized HBC similar to HBCA1. The block copolymer has a strong ability to form well-defined PLL-α-helices. The complex self-assembles in two Colh phases, depending on the temperature. Both Colh phases possess a 2D hexagonal sub-lattice from the lateral organization of HBC columns and a 2D hexagonal superstructure

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139

x = 0.936 y = 0.064

HBCA1

R=

HBCA2

Chart 11 Structures of columnar LC phase-forming HBC-polymer ionic complexes

pffiffiffi with a lattice constant that is 7 times of the former, due to having six HBC columns attached to each α-helix PLL by ionic bonds. Interestingly, the intracolumn stacking order of HBC in the lower-temperature (below 54  C) Colh phase is significantly poorer than that in the high-temperature phase (Thünemann et al. 2003).

Columnar Phase-Forming Polymers with Discotic Mesogens in the Main-Chain TP units have also been integrated into the main-chain of a polymer, usually in the form of polyester. With an appropriate set of spacer/side-chain, a columnar LC phase can form in main-chain polymers shown in Chart 12. A discotic mesophase has been suggested for the PTP1-5-m, the first reported main-chain TP polymers, on the basis of DSC and texture observations (Wenz 1985). With additional X-ray diffraction or 2 H NMR evidences, the Colh phase of PTP2-5-14 (Kreuder et al. 1985; HerrmannSchoenherr et al. 1986), PTP3 (Hüser et al. 1989), PTP4-5-12, PTP4-5-14, and PTP4-7-14 (Kranig et al. 1990) has been established. The increase of spacer length monotonically reduces the mesophase clearing point. PTP4-5-20 is an amorphous material over the entire studied temperature range (Kranig et al. 1990). The alignment behavior of columnar phase-forming TP main-chain polymers is quite different from their side-chain counterparts. In the Colh phase of a main-chain polymer, the column axis is aligned perpendicular to the strain direction, while the columns in side-chain polymers are parallel to the strain direction (Hüser et al. 1989). A columnar mesophase has also been suggested for PolyPc1, a polymer containing phthalocyanine units in the main-chain. This is in contrast to the absence of columnar phase-forming side-chain phthalocyanine polymers, at least to the best knowledge of the author. Although polyacrylate-, polymethacrylate (Kamachi et al. 1987)-, polystyrene-, and polysiloxane (Makhseed et al. 1999)-bearing phthalocyanine side-groups have been prepared, none of them form a columnar mesophase. Monomeric phthalocyanines, on the other hand, have a strong ability to form a columnar mesophase, when equipped with appropriate flexible side-chains.

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PTP1-5-m, m = 8, 10, 12, 14, 16

PTP3

PTP2-5-m, m = 8, 14, 20

PTP4-p-m, p =5, 7; p =12, 14, 20

PolyPc1, R = C8H17

Chart 12 Structures of columnar LC phase-forming polymers with mesogens in the main-chain

Cross-Linked Columnar LC Phase-Forming Polymers Columnar mesophases may also form in cross-linked polymers which are typically prepared by polymerizing a multifunctional monomer in its columnar LC phase. All three HBC monomers shown in Chart 13 exhibit a broad Colh phase temperature range suitable for thermal polymerization. The polymerization of HBCV1 gives a very low conversion and a product mixture with DP up to 4, probably due to the low concentration of polymerizable groups. In contrast, thermal polymerization of HBCV2A and HBCV2MA leads to corresponding cross-linked polymers with ~60% conversion. Both cross-linked polymers exhibit a highly ordered Colh phase in a very broad temperature range from 50 to 300  C (decomposition) (Brand et al. 2000). Similar to HBCs, phthalocyanines PcHA, PcHMA, and PcCuMA were successfully polymerized thermally in the Colh phase. X-ray diffraction has shown that the resulting cross-linked polymers retain the highly ordered Colh phase with a tight π-stack (intra-stack spacing ~ 0.33 nm) (van der Pol et al. 1990). Besides the addition reaction of vinyl groups, azide-alkyne [3 + 2] cycloaddition was also utilized to cross-link discotic molecules (Kayal et al. 2013). A 1:1 mixture of CuTAPA3 and CuTAPB3 exhibits its Colh phase over 50 to 100  C. The mixture was slowly cross-linked via thermally activated azide-alkyne [3 + 2] cycloaddition at 65  C while in its Colh phase. The X-ray diffraction results showed that the phase structure hardly changes during the reaction which eventually leads to an insoluble 3D network polymer in its Colh phase. In contrast, the cross-linking reaction of thesame mixture in solution, catalyzed by Cu+, produced an amorphous insoluble polymer, highlighting the importance of carrying out cross-link reaction with the monomer in its columnar LC phase.

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R HBCV1

M

HBCV2A

PcHA

HBCV2MA

PcHMA

HH

PcCuMA

Cu

R

R

HH

CuTAPA3 (CH2)3N3 CuTAPB3

Chart 13 Monomers of cross-linked columnar LC phase-forming polymers

Future Directions Although columnar phases are not always associated with discotics, discoticscontaining polymers may play a major role in future directions of columnar phaseforming polymers, largely thanks to their attractive potential applications in the field of optoelectronics, arising from π-stacked discotic cores. It has been shown that the attachment of triphenylene units to a macromolecule may give rise to a π-stacked columnar mesophase featuring an order (Mu et al. 2017) and/or molecular mobility (Kranig et al. 1990) supporting charge transport even better than the small molecule counterpart in a highly ordered columnar mesophase (Mu et al. 2017), which makes such polymers particularly appealing. Therefore, it is of great interest to extend such advantageous phase structures to those polymers containing discotic units that are more capable of supporting high carrier mobility. Considering that there are no reports on columnar phase-forming polymers covalently bearing hexabenzo[bc,ef, hi,kl,no,qr]coronene, phthalocyanine, or porphyrin as the side-groups, this necessitates a deeper understanding on the polymer structure-phase relationship.

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Fullerene Liquid Crystals Xiaoming Yang, Tiantian Zhu, and Yingfeng Tu

Contents Definition . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . Introduction . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . General Description of [60]Fullerenes . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . General Introduction of Fullerene Liquid Crystals . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . Design Strategies for Fullerene LCs . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . Molecular LC Approach . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . Supramolecular Approach . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . Properties of Fullerene LCs . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . Electrochemical Properties . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . Chiral Properties . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . Light Emitting Properties . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . Optoelectronic Properties . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . Challenges . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . References . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . .

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Abstract

Fullerene-based liquid crystalline materials contain both the excellent optical and electrical properties of fullerene and the self-organization and external-fieldresponsive properties of liquid crystals (LCs). They have potential applications in optical and photovoltaic devices, organic field-effect transistors, especially as active materials in polymer solar cells. Here, we have summarized the results on the design strategies for [60]fullerene LCs into two approaches, namely, the molecular LC approach and the supramolecular LC approach, respectively. The molecular LC approach is introduced first to design [60]fullerene LCs via a X. Yang · T. Zhu · Y. Tu (*) College of Chemistry, Chemical Engineering and Materials Science, Soochow University, Suzhou, China e-mail: [email protected]; [email protected]; [email protected] © Springer Nature Switzerland AG 2020 L. Zhu, C. Y. Li (eds.), Liquid Crystalline Polymers, Polymers and Polymeric Composites: A Reference Series, https://doi.org/10.1007/978-3-030-43350-5_62

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traditional way, where C60 was linked together with large liquid crystal mesogens to give nematic, cholesteric, smectic, and columnar phases of LCs by the Bingel reaction or the 1,3-dipolar cycloaddition reaction. This strategy needs large mesogens to fulfill the molecular aspect ratio of LCs. Thus the content of C60 is low, usually less than 20%, and the properties from fullerene are hindered. Another approach is via the supramolecular self-assembly, where the [60]fullerene derivatives self-assemble to form supramolecular structure which meets the aspect ratio requirements of liquid crystals. The supramolecular LC approach opens up new strategies to design liquid crystals with low molecular aspect ratio, enabling high [60]fullerene content in the supramolecular LCs. Some of the interesting properties form [60]fullerene LCs are also summarized, especially the electrochemical properties, the chiral properties, the light emitting properties, and the optoelectronic properties. Among them, high fullerene content is the key for these LCs to achieve good properties. Keywords

Fullerene · Liquid crystal · Supramolecular chemistry · Self-organization

Definition [60]Fullerene-based liquid crystals are a new kind of materials with excellent electro-optical properties of fullerenes and self-organization and processible properties of liquid crystals. The design of fullerene liquid crystals includes two ways: direct synthesis of molecular liquid crystals with the molecular aspect ratio meeting the requirement for liquid crystals and supramolecular liquid crystals with the supramolecular structure meeting the aspect ratio requirement of liquid crystals through self-assembly or self-organization of fullerene derivatives. This entry will focus on the design method of [60]fullerene liquid crystals from molecular structure aspect. Functional groups such as ferrocene, dendron, chiral, and phthalocyanine can be introduced into fullerene-based liquid crystals by careful molecular design, which expands the application field of fullerene-based liquid crystals.

Introduction General Description of [60]Fullerenes [60]Fullerene (C60), the most familiar member in fullerene’s family, is a truncated icosahedron (Ih) with 60 vertices and 32 faces (20 hexagons and 12 pentagons where no pentagons share a vertex) consisting of sp2-hybridized carbon atoms, as shown in Fig. 1. The van der Waals diameter of a C60 molecule is 1.01 nanometer (nm), while the nucleus to nucleus diameter of a C60 molecule is 0.71 nm. The C60 molecule has two bond lengths. The 6:6 ring bonds (between two hexagons) can be considered as

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Fig. 1 Molecular structure of C60

“double bonds” and are shorter than the 6:5 bonds (between a hexagon and a pentagon). Its average bond length is 0.14 nm. Each carbon atom in the structure is bonded covalently with three others. Detailed physical and chemical properties about fullerenes can be found in books (Andreas et al. 2005; Kadish and Ruoff 2000). C60 is the most extensively studied carbon-based materials in recent years due to the excellent redox, optical, and optoelectronic properties. The optical absorption properties of C60 derivatives match solar spectrum in a way that suggests that C60based films could be useful for photovoltaic applications. Because of its high electronic affinity, fullerene derivatives are one of the most common electron acceptors used in donor/acceptor-based solar cells. Conversion efficiencies up to 10.4% have been reported in C60-polymer cells (Liu et al. 2014). In addition, C60 is one kind of promising electron-transport material as an n-type semiconductor exhibiting relatively high carrier mobility of approximately 1.0 cm2 V
1 s
1 in the single crystal and polycrystalline thin films (Anthopoulos et al. 2006).

General Introduction of Fullerene Liquid Crystals Liquid crystals (LCs) are one type of materials that self-organize into complex, hierarchical structures. Functional liquid crystal assemblies offer the possibility of fabricating dynamic, addressable structures where the functional moieties are organized in a predetermined and controllable fashion. Fullerene-containing liquid crystals thus attracted the interest since they raise the hope that one can draw both the excellent opto-electrical properties (from fullerene) and the external-fieldresponsive properties (from LCs) out of such materials. In addition, fullerenecontaining liquid crystals would provide fundamental information for a better understanding of the factors which govern the formation of supramolecular structures obtained from the organization of fullerene-containing molecules. However, it is not easy to construct [60]fullerene LCs by traditional way. Covalent attachment of fullerene molecules to a mesomorphic promoter molecule such as cyanobiphenyl, cholesterol, and dendrimers did not result in the desired LC properties in some cases, because the grafted fullerene groups tend to disturb the

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organization of the original mesogenic molecules (Tirelli et al. 2000). This is due to the big spherical structure of C60 which is not easy to construct molecules with aspect ratio >4. In fact, fullerenes belong to the family of plastic crystals, which is in contrary to LCs by definition (André et al. 1992; Heiney 1992). A plastic crystal is a crystal composed of weakly interacting molecules that possess some orientational or conformational degree of freedom, i.e., has positional order but no orientation order, while liquid crystals have orientation order and no long-range positional order. The shape structures of plastic crystal and liquid crystal are totally different: for liquid crystals, they have anisotropic shape with rod, disk, or bowl-like structure, while for plastic crystals, the shape is highly symmetric (usually sphere-like). The question then rise: can we construct LCs from typical plastic crystal moiety like fullerene? To now, there are typically two strategies for the design of [60]fullerene LCs, namely, the molecular LC approach and the supramolecular LC approach. The molecular LC approach has two different ways to meet the aspect ratio requirement of LC: (1) by connecting fullerenes with a very long mesogen or (2) by multiaddition of mesogens on fullerene cage. In the following, we will first discuss the synthesis of [60]fullerene LCs based on the above design strategy, then the properties from these [60]fullerene LCs, and last our perspectives for these materials.

Design Strategies for Fullerene LCs Molecular LC Approach Taking into account the size and shape of the C60 unite, the elaboration of mesomorphic fullerene derivatives represents a conceptual and synthetic challenge. [60] Fullerene (C60) was successfully incorporated into liquid crystal systems (Chuard and Deschenaux 1996; Dardel et al. 1999; Tirelli et al. 2000; Felder-Flesch et al. 2000, 2006; Even et al. 2001; Dardel et al. 2001; Sawamura et al. 2002; Kimura et al. 2002; Campidelli et al. 2002, 2003, 2004, 2005, 2006a, b; Matsuo et al. 2004; Allard et al. 2005; Bushby et al. 2005; Lenoble et al. 2006) to give nematic, chiral nematic, smectic A, smectic B, columnar, and cubic phases. Recently, it was shown that hexaaddition of a malonate derivative, which possessed a monotropic nematic phase, yielded a fullerene derivative, which exhibited an enantiotropic smectic A phase. This result indicated that C60 can be used as a platform for the synthesis of liquid crystalline materials from a non-mesomorphic addend. A bischolesterol derivative was used to prepare a hexa-adduct of C60. As expected (due to the cholesterol units), a smectic A phase was obtained.

Linking Long Mesogens with Fullerene The first [60]fullerene-containing thermotropic liquid crystal 1 was synthesized in 1996 (Fig. 2). The following structural requirements were applied for the successful design of the first [60]fullerene-containing LC derivative: (1) to generate strong intermolecular interactions between the mesogenic units, a twin cholesterol

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Fig. 2 Molecular structure of the first fullerene-containing liquid crystal (1) and fullereneferrocene compound (2, 3)

framework was selected for the formation of C60 derivative; (2) to lower the transition temperatures, a flexible chain was used as a spacer between the cholesterol derivative and the C60 moiety; and (3) owing to the well-established synthetic procedure, the formation of a methanofullerene was chosen to connect the cholesterol fragment to the C60 (Chuard and Deschenaux 1996). The targeted fullerene derivative 1 showed a monotropic smectic A phase. Following this strategy, Deschenaux et al. introduced ferrocene and fullerene together in the LC, with some of the structure presented in Fig. 2 (2 and 3) (Deschenaux et al. 1998). The design of liquid crystalline fullerene-ferrocene compounds is attractive since such structures lead to new multicomponent mesomorphic materials. The mixed fullerene-ferrocene thermotropic liquid crystal 2 represents the first member of this new family of mesomorphic materials. More fullerene-

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containing thermotropic liquid crystals were obtained by addition of mesomorphic malonates to C60 following the Bingel reaction (Chuard et al. 2000). The smectic A phase was observed in fullerene derivative 3.

Multiaddition of Mesogens on Fullerene A hexa-adduct 4 and mono-adduct 5 of [60]fullerene (Fig. 3) were synthesized by addition of a mesomorphic twin cyanobiphenyl malonate derivative to C60 (Chuard et al. 1999). For mono-adduct 5, the intermolecular interactions are too weak to generate mesomorphism, while hexa-adduct 4 forms focal-conic and homeotropic textures, the typical textures of smectic A phase. It is noteworthy that polyaddition can be used for the preparation of fullerenecontaining thermotropic liquid crystals from addends with low liquid crystalline tendencies. Highly symmetric methanofullerene adducts, bearing long alkyl chains without (6) or with (7, 8, and 9) mesogenic groups (Fig. 4), were synthesized (Tirelli et al. 2000). The novel fullerene adducts 8 and 9 exhibit mesogenic properties, while 6 and 7 do not. The formation of liquid crystalline phases could only be achieved with well-known mesogens and an appropriate balance of their number per fullerene unit. In order to explore the influence of polyaddition on the formation of chiral liquid crystalline phases, mono-(10) and hexa-(11) adducts of C60 were synthetized (Fig. 5) (Campidelli et al. 2006a). The latter structure is interesting for the exploration of the material properties of the chiral mesogens when assembled around a focal point. Hexa-adduct 11 displayed a schlieren and pseudo-focal conic textures and exhibited a chiral nematic phase. A value of 2.0 μm for the pitch length was obtained at room temperature. The fullerene moiety in 11 is shielded very effectively among the laterally attached mesogens, without disturbing the helical supramolecular organization of the mesophase. C60 is better embedded within the self-organizing system in the hexa-adduct than in the dendritic mono-adduct, because the addends are symmetrically distributed all over the fullerene sphere effectively, whereas in the mono-adduct 10, C60 is left comparatively more exposed, increasing the possibility of aggregation interactions between the C60 cages, which are detrimental to mesophase formation.

Fig. 3 Hexa-adduct 4 and mono-adduct 5

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Fig. 4 Molecular structures of C60 adducts investigated by multiaddition of mesogen on fullerenes (6–9)

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Fig. 5 Structure of mono-adduct (10) and hexa-adduct (11) of fullerene

As a general indication, the C60 structure was shown neither to directly exhibit a mesogenic order nor to enhance the one coming from otherwise active mesogens. In case of liquid crystals, C60 is hidden in the mesogenic cloud and does not alter the supramolecular organization encoded in the mesogens. Actually, C60 unit tends to lower the stabilization of the liquid crystalline phase, probably because of steric effects arising from its bulky spherical shape. Thus, the spherical shape of C60 is not conducive toward the formation of calamitic liquid crystal phases, and it acts like a nonmesogenic dopant. The content of fullerene in multi-adduct is very low, and the fullerenes’ supreme opto-electrical properties are hindered.

Supramolecular Approach Fullerene LCs with Columnar Supramolecular Structure Previous fullerene-containing LCs were built by connecting fullerene molecules to an inherently mesogenic unit. In 2002, with the development of supramolecular LCs, Nakamura et al. reported the first supramolecular [60]fullerene LC (Sawamura et al. 2002). The chemical structure of the supramolecular fullerene LCs (12–14) is shown in Fig. 6, where the C60 derivatives self-assemble to form anisotropic “nano-shuttlecocks” that stack into a one-dimensional columnar supramolecular structure. This design strategy

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Fig. 6 Molecular structure of the series of 12–14 and top view, side view, and a stack of five molecules of the series of compound with the alkyl chain length n = 12. Color code: red, fullerene core; blue, aromatic groups; and gray, alkyl chains. (Reprinted by permission from (Sawamura et al. 2002), copyright 2002 Macmillan Publishers Ltd.)

should be applicable to a range of other molecules and materials and opens a facile way to build a variety of fullerene LCs since the molecular aspect ratio is not needed by this strategy. In the early design (Fig. 6), the five aryl groups (R) form a cone shape, protruding directly from the fullerene core and ensuring tight fitting of another fullerene molecule in the cone to form a rigid column in crystals and LCs. The tightness of this connection restricts the structural mobility of the column. The stacking is driven by attractive interactions between the spherical fullerene moiety and the hollow cone formed by the five aromatic side groups of a neighboring molecule in the same column. Following this strategy, the fullerene supramolecular LCs with a larger cavity that can comfortably accommodate the second fullerene molecule and allows a more structurally flexible connection between the molecules in the column were also reported by them (Matsuo et al. 2004, 2006). By introducing ferrocene group in the structure, the fullerene LC showed reversible multi-electron redox behavior, accepting and giving up a total of at least four electrons. In addition, by introducing a conical shape, a polar iron-ferrocene complex, and long alkyl chains into a fullerene molecule, the fullerene dipolar molecules 15 and 16 (Fig. 7) underwent the microphase separation and formed the three-dimensional lattice in a crystalline and a thermotropic liquid crystalline phase. The key feature is a tetrameric octupole-like aggregate, in which four dipoles are arranged in a supramolecular way to cancel the molecular polarity, forming a sphere (Li et al. 2010).

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Fig. 7 A schematic illustration of dipolar conical molecules 15 and 16

Fullerene LCs with Lamellar Supramolecular Structure When the head-to-tail stacking is prohibited, lamella stacking will be preferred, as reported by Matsuo and Nakamura et al. in their [60]fullerene derivatives 17–22 (Fig. 8). The fullerene core is connected to many hydrocarbon groups by silyl acetylene tethering (Zhong et al. 2007), where the silicon atom facilitates installation of the hydrocarbon substituents through a simple Grignard reaction. A methyl group was installed to the center of the molecular cavity to prevent the head-to-tail stacking. The lamellar structure can be formed by the crystal 17 and the liquid crystal 20 where the distance between the fullerene layers is 22.44 Å and 22.6 Å separately. The molecules are arranged alternately with the R groups and the methyl group attached to the fullerene core upward and downward. The short distances (9.80 and 10.16 Å) between neighboring fullerene cores suggest a strong fullerene/fullerene interaction within the same layer. In 2008, Li et al. reported another fullerene LC system with lamellar assembly structure from amphiphilic oligothiophene-C60 dyad (Fig. 9). The molecular structures of 23 and 24 are similar to each other except for the terminal wedges (Li et al. 2008). Compound 23 bears a hydrophilic wedge with triethylene glycol chains and, on the other side, a hydrophobic wedge with paraffinic chains. Dyad 23 formed a LC smectic A mesophase over a wide temperature range from 136.1  C to 18.3  C. Polarized optical microscopy (POM) of LC 23 displayed a typical focal conic texture. Synchrotron radiation small-angle X-ray scattering (SAXS) analysis showed sharp peaks with dspacings of 10.6, 5.3, 3.5, and 2.6 nm, which can be indexed as (100), (200), (300), and (400) reflections of a lamellar structure with a layer width of 10.6 nm.

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Fig. 8 Molecular structure of crystals of 17 and 18 and liquid crystals of 19–22

Fig. 9 Structure of oligothiophene-containing [60]fullerene LC 23 and 24

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Fig. 10 (Left) Structure of C60 derivatives (25–30); (Right) illustration for the proposed lamellar organization of 25. (Reprinted with permission from (Nakanishi et al. 2008). Copyright (2008) American Chemical Society)

Similar to 23, nonamphiphilic 24 formed a smectic A LC mesophase in a temperature range from 111.4  C to 12.5  C. SAXS pattern displayed intense and weak diffractions with d-spacing of 5.7 and 2.9 nm, respectively, which were indexed as (100) and (200) reflections of a lamellar structure with a layer width of 5.7 nm. For the abovementioned C60-containing liquid crystals, the C60 content (based on molecular weight) in liquid crystals is usually ~20% because of their bulky appendage (Chuard and Deschenaux 1996, 2002; Sawamura et al. 2002; Bushby et al. 2005; Felder et al. 2000; Zhong et al. 2007; Deschenaux et al. 2007). The low content of fullerenes hinders their excellent optoelectronic properties. How to increase the content of C60 while maintaining the LC properties is a main challenge for the C60-containing LCs, since the low content of C60 will hinder their application. The breakthrough was made by Nakanishi et al., where they synthesized multi (alkyloxy)phenyl group functionalized with fulleropyrrolidines (Fig. 10), which show LC properties with a high C60 content (up to 50%). The simple modification of C60 by two (26) or three long alkyl chains (25 and 28) permits a high C60 content in the mesomorphic materials. Supramolecular fullerene nano- and microarchitectures were formed via the intermolecular phase separation introduced by C60 (π-π) and alkyl chain interactions (van der Waals). However, when further reducing the content of alkyl chains to improve the fullerene content, the approach is failed as compounds 27, 29, and 30 showed no LC properties. With such high C60 content, a comparably high electron mobility of 3  10
3 cm2 V
1 s
1 as well as electrochemical properties is observed (Nakanishi et al. 2008). Bent-shaped molecules, through dense arrangements that hinder the molecular rotation, lead to strong polar order within either layers or columns, in many cases producing supramolecular chirality by achiral molecules. The molecules can often be switched in these mesophases by external stimuli, and they have afforded exceptionally good macroscopic polarization values and piezoelectric, flexoelectric, and nonlinear optical responses. A new class of nondendritic hybrid C60-containing bent-shaped molecules (31–34) were reported by Vergara et al. (2011), as presented in Fig. 11.

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Fig. 11 Chemical structures of the C60-based bent-core compounds (31–34) prepared with LC properties

In the mesophase, each molecule is folded with the two bent-core units oriented in the same direction. The fullerenes occupy the central area of the bilayer in a head-tohead arrangement, and the hydrocarbon chains form an aliphatic region between neighboring layers.

Fullerene LCs with Helical Supramolecular Structure Studies on C60-based LC materials with a high fullerene content and with simpler molecular structures are scarce (Campidelli et al. 2003, 2005, 2006a, 2010; Nakanishi et al. 2008; Mamlouk et al. 2007; Perez et al. 2008; Orlandi et al. 2009; Fernandes et al. 2010), possibly because the appropriate combination of C60 units with other organic functional structures to attain mesomorphism is not an easy task. Hayashi et al. reported the synthesis of LC and charge-transport properties of a zinc phthalocyanine (ZnPc)-C60 dyad 35, as presented in Fig. 12 (Hayashi et al. 2011). Six 4-dodecyloxyphenoxy groups were introduced into the periphery of the central core to form a discotic columnar structure of the ZnPc. C60 was also tethered to the ZnPc core via a short, semiflexible bridge. Due to the strong π-π interaction between the C60 molecules and the covalent linkage, the C60 molecules would be arranged successively along the ZnPc 1D column, leading to D-A bicontinuous structure in the LCs. Wang et al. reported the design and synthesis of porphyrin-C60 dyad (36 in Fig. 13) and investigated the supramolecular structure, where geometrically isotropic

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Fig. 12 Phthalocyanine fullerene dyad 35 and the self-assembled supramolecular helical columnar structure. (Reprinted with permission from (Hayashi et al. 2011). Copyright (2011) American Chemical Society)

(C60) and anisotropic (porphyrin) units coexist (Wang et al. 2012). Structural analysis revealed that this stable phase possessed a supramolecular “double-cable” structure with one p-type porphyrin core columnar channel and three helical n-type C60 peripheral channels. These “double-cable” columns further packed into a hexagonal lattice with a = b = 4.65 nm, c = 41.3 nm, α = β = 90 , and γ = 120 . The column repeat unit was determined to possess a 12944 helix. With both donor (D, porphyrin) and acceptor (A, C60) units having their own connecting channels as well as the large D/A interface within the supramolecular “double-cable” structure, dyad 36 showed good photogenerated carriers with longer lifetimes compared to the conventional electron acceptor [6,6]-phenyl-C61-butyric acid methyl ester (PCBM). Star-shaped oligophenylenevinylene (OPV) mesogens are shape-persistent and possess formally large void space. A mesogen with three phenylacetylene repeating units packs densely in a columnar helical arrangement (Lehmann and Hugel 2015). Attachment of one fullerene through a short spacer results in an exceptional increase of the mesophase stability (38 and 39 in Fig. 14). The shape-persistent OPV stars maintain their overall star morphology and exhibit large void space between their arms. In the OPV-star 37 and star-OPV-fullerene hybrid 38, the void is perfectly filled by non-mesomorphic fullerene guests. These guests facilitate the space filling and stacking in columnar structures and consequently increase the mesophase stability by more than 70  C compared to the parent mesogen 37. Fullerene is a spherical building block and has, among others, been incorporated in LC phases by coupling to mesogenic dendrons of various generations. In the hybrid derivative 39 and the mixture between 39 and 37, fullerene fills a great portion of the void space between the arms and facilitates the columnar stacking. The systems optimize the packing of fullerenes by the arrangement in a triple helix. Such highly ordered, columnar donor (stilbenoid scaffold)-acceptor

Fig. 13 Chemical structure of porphyrin-C60 dyad 36 and illustration of the supramolecular “double-cable” helical structure and the porphyrin core channel and C60 peripheral channels in the 12944 helical structure. Alkyl tails are omitted for clarity. (Reprinted with permission from (Wang et al. 2012). Copyright (2012) John Wiley and Sons)

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Fig. 14 OPV stars (37) and fullerene-OPV hybrid stars (38 and 39)

(fullerene) assemblies are interesting materials for organic electronics and especially photovoltaic applications if they can be oriented at surfaces.

Lamella Supramolecular Structure with Sandwiched 2D Crystals As pointed above, the π-π interactions, which are considered as weak intermolecular forces comparable to hydrogen bond interactions, can promote the formation of fullerene LCs via supramolecular assembly. Very recently, we found that in a properly designed system of fullerene dyads, the fullerenes can self-organize to form two-dimensional (2D) crystals sandwiched within the flexible alkyl chain layers. The structure of the [60]fullerene derivatives 40–43 is shown in Fig. 15. A typical molecule from these dyads consists of a rigid [60]fullerene, a gallic ester segment substituted with three long alkyl chains as the soft part, and a multimethylene unit as a flexible spacer connecting the two segments (Zhang et al. 2015). The design strategy is to introduce a soft group onto fullerene (the rigid component), so the molecules will self-organize to form supramolecules as a result of π-π interactions between fullerenes and phase separation between the rigid and the soft parts. The [60]fullerene dyads undergo self-organization by π-π interactions between fullerenes to form triple-layer 2D crystals, with the soft alkyl chains forming

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Fig. 15 Molecular structure of fullerene dyads 40–43 (left), illustration of the sandwiched 2D crystals formed by the self-organization via π-π interaction and phase separation (middle), and the supramolecular liquid crystals formed by the random stacking of 2D crystals (right). (Reprinted with permission from (Zhang et al. 2015). Copyright (2015) John Wiley and Sons)

continuous lamellar layers above and below the 2D crystals, as shown in Fig. 15 (middle). The random packing of 2D crystal-containing lamellar gives rise to the formation of a new family of thermotropic supramolecular LCs. It should be noted that the flexible spacer plays an important role to decouple the interactions between the soft part and the rigid part. When the spacer is very short, no LC mesophase was observed from 43, while those with long spacer (40–42) showed LC properties. This design strategy should be applicable to other molecules and lead to an enlarged family of 2D crystals and supramolecular liquid crystals. As fullerenes are part of the LC mesogen, these supramolecular LCs have high fullerene content.

Properties of Fullerene LCs The successful creation of fullerene-containing LCs opens avenues to the design of new and unique mesomorphic functional materials. Such materials combine the exceptional electrochemical and photophysical properties of fullerenes with those of self-organizing and self-assembling media. The association of such properties could be used in the evolution of solar cell technologies and in the development of supramolecular redox and optical molecular switches based on photoinduced electron transfer, for example, between ferrocene and fullerene units.

Electrochemical Properties Fullerene (C60) is a unique electroactive material; it can reversibly accept up to six electrons. Therefore, incorporation of C60 into liquid crystals offers the possibility of combining the efficient acceptor characteristics of C60 with the anisotropy of liquid crystals. Mixed fullerene-ferrocene liquid crystals are interesting materials from the

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Fig. 16 Structures of the liquid crystalline fullerene-ferrocene dyad 44

point of view of electrochemical properties, since fullerene is acted as electron acceptor subunits while ferrocene as an electron donor subunit (Fig. 2). The electrochemical behavior of the fullerene-ferrocene dyads has been investigated (Carano et al. 2002). Both fullerene- and ferrocene-centered oxidation and reduction processes were observed in fullerene-ferrocene dyads: a total of seven reduction peaks were detected in the negative potential region and two oxidation peaks in the positive potential one. The electrochemical property of fullerene-peralkylated ferrocene derivative was investigated, where its chemical structure is shown in Fig. 16. Compound 44 contains a second-generation liquid crystalline dendrimer which provides the mesomorphic properties (Campidelli et al. 2008). The liquid crystalline dyad 44 displayed an enantiotropic smectic A phase from 57  C to 155  C. The dyad showed electrochemical activity when measured as cast films on a glassy carbon electrode above the solid/mesomorphic phase transition temperatures. A cast film of 44 in 0.1 M aqueous n-Bu4NCl solution at 60  C showed the first and second redox events corresponding to the generation of C60 monoanion and dianion at potentials of Ered,1 at
0.70 V and Ered, 2 at
0.87 V, respectively.

Chiral Properties Chiral nematic liquid crystals form helical superstructures where the local director of the nematic phase is rotated in either right- or left-handed direction perpendicular to the molecular long axes. To induce chirality into C60 systems, laterally branched chiral mesogenic groups were introduced, within the malonate addend (Campidelli et al. 2003), as shown in Fig. 17 (45).

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Fig. 17 Chemical structure of fullerene LC 45 that forms chiral nematic phase

Light Emitting Properties Metal-containing fullerenes, and particularly trimetallic nitride template endohedral metallofullerenes (TNTEMFs), elicit increasing attention for their outstanding electronic and optical properties. Light emission efficiencies are closely related to the molecular organization and degree of ordering. It is widely recognized that selforganization by the formation of low-dimensional LC phases is a key strategy for controlling the ordering and structuring of organic semiconductors because it helps to reduce or even suppress defect formation. The grafting of the mesomorphic double oligo(phenylene ethynylene) (dOPE) onto the fullerenes promotes mesomorphism in the fullerene adduct 46, with the induction of columnar phases resulting from the triple segregation between the fullerene (core), mesogens (walls), and chains (continuous medium) according to a Kagome lattice (Fig. 18). The liquid crystalline derivative of Y3N@C60 shows remarkable photophysical properties: OPE units act as 100% efficient lightharvesting antennae to sensitize a bright and long-lived fullerene core emission. These luminescence properties are retained in the mesophase and, coupled to a quite strong absorption in the visible region extending up to 750 nm, open up a variety of potential applications.

Optoelectronic Properties The combination of electron acceptor (fullerene) and electron donor (ferrocene, phthalocyanines) within the same architecture opens the door for the construction of photoactive molecular devices. A simultaneous application of a phthalocyanine and a fullerene to LCs is highly attractive to form such D-A heterojunction structures. Liquid crystalline donor (i.e., phthalocyanine) was covalently linked to fullerene to achieve efficient charge-transport properties in a liquid crystalline phase. The columnar structure exhibited highly efficient ambipolar charge-transport character, demonstrating the potential utility of the strategy in organic electronics. The D-A heterojunction structure of ZnPc-C60 dyad 35 was found to exhibit highly efficient ambipolar charge-transport properties (Hayashi et al. 2011) (Fig. 12). ZnPc-C60 displays remarkably high hole mobility of μh = 0.26 cm2 V
1 s
1 and

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Fig. 18 Chemical structure of Y3N@C60-based dyad 46

electron mobility of μe = 0.11 cm2 V
1 s
1, which are the highest values reported for organic materials with D-A heterojunction. The fullerene-ferrocene derivative (Fig. 2) was also designed to develop photoactive liquid crystal switches by combining two sets of data: firstly, electron transfer was used to generate liquid crystalline ferrocenium derivatives from non-mesomorphic ferrocenes and, secondly, photoinduced electron transfer from ferrocene to C60 was shown to occur in fullerene-ferrocene dyads. Cholesterol was used as liquid crystalline promoter. X-ray diffraction experiments and volumetric measurements indicated that the fullerene-ferrocene derivative is organized in double-layered structures. On the other hand, photoinduced electron transfer in fullerene-ferrocene liquid crystals could be used to control the liquid crystalline properties because of the presence of either the ferrocene (light off) or ferrocenium (light on) species (Deschenaux et al. 1998) (Fig. 2). Photophysical studies revealed that electron transfer occurs from the donor ferrocene to the electron accepting fullerene. The formation of a long-lived radical pair, with lifetimes of the order of several hundred nanoseconds, was confirmed (Even et al. 2001).

Challenges To now, most of the research works are focused on the synthesis of new [60]fullerene liquid crystals and the characterization of their phase diagram and hierarchical structures, while less of their properties have been reported. To achieve good properties, the high fullerene content in liquid crystals is the key to display fullerene’s properties. The supramolecular approach seems more promising for this goal, where functional groups can be introduced in while less molecular aspect ratio is concerned. The challenge here is to make a balance between the content of introduced functional groups to achieve liquid crystalline properties and the content of

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fullerenes to achieve their good optoelectronic properties. On the other hand, how to use the fullerene liquid crystals’ external responsibility to achieve better optoelectrical properties is an important yet currently untouched field. This leads to an interesting nanotechnology, i.e., arrange [60]fullerene cages in a desired way to achieve nanolithography.

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Structure and Assembly of Liquid Crystalline Block Copolymers Kishore K. Tenneti, Xiaofang Chen, Qiwei Pan, and Christopher Y. Li

Contents Definition . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . Introduction . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . Phase Behavior of Block Copolymers . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . Liquid Crystals and Liquid Crystalline Polymers . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . Liquid Crystals . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . Liquid Crystalline Polymers . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . Liquid Crystalline Block Copolymers . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . LCBCPs with Covalent Interactions: RCBCPs . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . LCBCPs with Covalent Interactions: Mesogen Jacketed Liquid Crystalline Block Copolymers . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . LCBCPs with Covalent Interactions: Side-Chain Liquid Crystalline Block Copolymers (SC-LCBCPs) . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . Phase Behavior of LCBCPs with Noncovalent Interactions: H-Bonded Liquid Crystalline Block Copolymers . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . Conclusion and Outlooks . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . References . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . .

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K. K. Tenneti · C. Y. Li (*) Department of Materials Science and Engineering, Drexel University, Philadelphia, PA, USA e-mail: [email protected]; [email protected] X. Chen Department of Materials Science and Engineering, Drexel University, Philadelphia, PA, USA Suzhou Key Laboratory of Macromolecular Design and Precision Synthesis, Jiangsu Key Laboratory of Advanced Functional Polymer Design and Application, College of Chemistry, Chemical Engineering and Materials Science, Soochow University, Suzhou, People’s Republic of China e-mail: [email protected] Q. Pan Department of Materials Science and Engineering, Drexel University, Philadelphia, PA, USA Department of Materials Science and Engineering, South China University of Technology, Guangzhou, China e-mail: [email protected] © Springer Nature Switzerland AG 2020 L. Zhu, C. Y. Li (eds.), Liquid Crystalline Polymers, Polymers and Polymeric Composites: A Reference Series, https://doi.org/10.1007/978-3-030-43350-5_64

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Abstract

Liquid crystalline block copolymers have been extensively investigated in the past two decades. Liquid crystals and block copolymers possess different ordering length scales. When they are chemically or physically incorporated in one materials systems such as liquid crystalline block copolymers, hierarchically ordered structures can be achieved with the ordering scales ranging from a few to a few tens nanometers. In this chapter, we will focus on the structure and assembly of such hierarchical systems in bulk states and the competition of the two different ordering processes. Keywords

Liquid crystals · Liquid crystalline polymers · Block copolymers · Self assembly

Definition Liquid crystalline block copolymers are block copolymers that are comprised of multiblocks and at least one of the blocks is formed by liquid crystalline polymers. This chapter reviews the structure and assembly behavior of liquid crystalline block copolymers.

Introduction Dealing with the scale of one to a hundred nanometers, nanostructured materials, and nanotechnologies have become an extremely active and vital area of research extending into almost every field of science and engineering. Materials with the size of this length scale show novel and exciting properties as predicted by Richard Feynman 43 years ago. There are two categories of methods to manufacture nanostructured materials: the so-called top-down and bottom-up methods. The top-down method typically begins with a suitable starting material and then “sculpts” functionality from it. On the other hand, the bottom-up approach first forms the nanostructured building blocks, and then assembles them into the final materials. As the top-down method, such as optical lithography technique, reaches its theoretic limits and processes such as the e-beam lithography are extremely costly and inefficient, bottom-up methods have drawn more and more attention in the scientific community. One key bottom-up approach is self-assembly, which is defined as the autonomous organization of components into patterns or structures without human intervention. Self-assembly occurs at all scales ranging from nanometer (selfassembled monolayers, liquid crystals, etc.), micrometer (colloidal crystals, etc.) to even macroscale (bacterial colonies, fish schools, etc.). It is believed that selfassembly on the nanometer scale is one of the few practical strategies for making ensembles of nanostructures.

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Self-assembly that occurs at more than one length scale is known as hierarchical self-assembly. From the technological point of view, hierarchical structures are of particular interest since they can blend the complexity into the structure and lead to multiple functional systems. One of the critical challenges that nanotechnology faces is the transfer of the novel properties obtained at nanoscale into a higher length scale for application reasons. Well-correlated structure at different length scales offers one of the ideal strategies to fulfill this goal. Due to their long-chain-nature, polymeric materials are known to be suitable for creating hierarchical self-assembly systems. As a typical biopolymer, natural protein possesses a specific amino acid sequence (primary structure), leading to its unique secondary, tertiary, and quaternary structures occurring at different length scales. These precise structure orderings also lead to their specific function(s). Compared to the precision and elegant complexity of biological systems, synthetic polymers are, in contrast, known as possessing relatively simple structures. However, linking different functional groups into synthetic polymer backbones may lead to the synthetic hierarchical structures possessing multiple functionalities that can mimic the complicated biopolymer structures and functions. To design a self-assembled molecular system with hierarchical structures, one needs to achieve structural ordering at different length scales. Structural ordering processes result from the competing molecular interactions (e.g., interactions between hydrophobic versus hydrophilic components, van der Waals, Coulombic, and hydrogen bonding). A number of different hierarchical assembly systems have been successfully developed by employing one or more of these interactions. Liquid crystalline (LC) block copolymers (BCPs) play a major role in creating hierarchical self-assembled structures. LC molecules typically consist of an anisotropic aromatic mesogen (with rod or disc shape) and aliphatic tails. In this chapter, we shall discuss the hierarchical structure and assembly behaviors of LC BCPs in bulk states. We shall first introduce BCP, followed by LC and LCPs. We will then summarize the hierarchical structure of LC BCPs.

Phase Behavior of Block Copolymers A BCP is a single phase macromolecular system formed by covalently combining two or more polymers. The enthalpic and entropic factors associated with covalently linking dissimilar polymers leads to intriguing phase behavior in BCPs (Bates 1991; Bates and Fredrickson 1990; Bates and Fredrickson 1999; Hamley 1998; Muthukumar et al. 1997). The thermodynamic interaction between two dissimilar molecules A and B is given by the Flory-Huggins interaction parameter, χAB, which is the enthalpy term. The magnitude of χAB for two polymers is usually positive indicating that their mixing is not favored. In a typical blend of two polymers A and B, this chemical incompatibility drives the system to undergo macrophase separation leading to A-rich and B-rich domains, having characteristic dimensions in the micron length scale, such that there is minimal segment-segment contact at the interface. In an AB BCP, the chemical bond between the blocks prevents such

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macrophase separation. However, BCPs do undergo microphase separation and form a variety of self-assembled ordered nanostructures at much smaller length scales (typically 10-100 nm). The nanostructure formation process is the cumulative result of enthalpy gain due to polymer segregation (into domains) and the corresponding loss in entropy due to chain stretching away from the interface. The chemical bond between the two blocks becomes the boundary between the two nanodomains. The dimensions of these nanodomains are a function of the radius of gyration (Rg) of the polymer chains (which in turn is a function of degree of polymerization, N). The product χN of a BCP system determines the strength or degree of segregation of these nanoscale domains. Theoretical calculations of coil-coil BCPs show that if χN > > 10.5, BCP undergoes microphase separation and forms ordered structures separated by sharp phase boundaries, and the system is considered to be in the strong segregation limit (SSL). For χN < < 10.5, entropic factors dominate and the BCP system is in a state of disorder or in the weak segregation limit (WSL) (Bates and Fredrickson 1990; Hamley 1998; Muthukumar et al. 1997). While χN determines the strength of microphase separation, the morphology of the BCP nanostructure is determined by the volume fraction ( f ) of the combining blocks. 1-dimensional (1D) structures called lamellar (L) morphology is commonly observed at symmetric volume fractions ( fA ~ 0.5) in which alternating layers of polymer A and B are stacked together. Asymmetry in f leads to nonplanar morphologies with curved interfaces so that the stretching penalty of the majority block is reduced. The degree of curvature increases with increase in the degree of asymmetry leading to morphologies such as 3D honeycomb-like structure called the double gyroid (G) ( fA ~ 0.28–0.34), 2D hexagonally close packed cylinders (C) ( fA ~ 0.17–0.28), and 0D spheres (S) ( fA < 0.17) of the minority block forming a body-centered cubic (BCC) lattice in the matrix of the majority block. L, G, C, and S, schematically shown in Fig. 1 (Grason 2006), are the four traditional BCP morphologies and are shown to be the only structures that possess a constant mean curvature at the corresponding f and are hence thermodynamically stable. A periodically ordered BCP phase structure turns into a disordered melt as a function of temperature. The temperature at which the order to disorder transition occurs is called order-disorder transition temperature (ODT), and it varies form system to system. Numerous research groups have investigated the phase behavior of coil-coil BCPs using both theoretical calculations and experimental methods. At the ODT, all the BCP phases (except the S morphology) undergo direct transition into a disordered state. The BCC S structure transforms into a disordered micellar structure which ultimately forms a disordered melt at two different temperatures called the lattice disorder temperature (LDOT) and demicellization temperature (DMT), respectively (Han et al. 2000). Helfandet al. and Semenov et al. provided the basis for quantitative analysis of BCP phase behavior in the SSL using a self-consistent field theory (SCFT) approach based on Meier’s criterion of enthalpic gain and entropic loss in high Mn systems (Helfand and Wasserman 1976; Helfand and Wasserman 1978; Meier 1969; Semenov 1985). Leibleret al. investigated the phase behavior of BCPs in the high temperature or WSL regions and in the regions near ODT (Leibler 1980). Matsenet al. developed a comprehensive theoretical phase

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Fig. 1 Traditional BCP morphologies (Grason 2006)

diagram using self-consistent mean-field approach that unified the SSL and WSL along with the intermediate segregation limit (ISL) and was recently modified as shown in Fig. 2a (Cochran et al. 2006; Matsen and Bates 1996). The phase diagram consists of the traditional BCP L, G, C, and S (both BCC and close packed) phases as a function of f. Advances in polymerization techniques (especially anionic) facilitated the synthesis of a variety of BCPs with very narrow Mn distributions (PDI = Mw/Mn, where Mw is the weight average molecular weight). Figure 2b shows the phase diagram obtained from experimental observations of phase structures in poly(styrene-bisoprene) (PS-b-PI) BCP system (Khandpur et al. 1995). The similarity between the two phase diagrams indicates the good agreement between theory and experimental observations. The phase behavior of PS-b-PI BCP system is one of the most extensively investigated (Bates et al. 1994; Forster et al. 1994; Hajduk et al. 1994, 1998; Hasegawa et al. 1987; Hashimoto et al. 1990; Khandpur et al. 1995; Winey et al. 1993a, b). Conventional L, G, C, and S phase structures where PS forms both minority and majority phase have been reported in samples with narrow PDI. In addition to these conventional BCP morphologies, complex morphologies such as perforated layered (PL) or catenoid structures were also observed where the majority phase forms regular perforations/channels in the domains of the minority blocks and these perforations possessed an in-plane hexagonal symmetry inside the minority domain. The PL structure is considered as a long-lived metastable state. Similar traditional BCP morphologies were also observed in a number of amorphous coil-coil BCP systems consisting of both low and high glass transition temperature (Tg) polymers such as poly(styrene-b-butadiene) (PS-b-PB) (Bates et al. 1982; Winey et al. 1991, 1992a, b), poly(styrene-b-methyl methacrylate) (PS-b-PMMA) (Akiyama and Jamieson 1992; Lowenhaupt and Hellmann 1991), poly(styrene-b-4-vinylpyridine) (PS-b-P4VP),(Saito et al. 1994; Schulz et al. 1996)

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Fig. 2 BCP phase diagram developed using (a) SCFT calculations (Cochran et al. 2006) and (b) experimental results from PS-b-PI system (Khandpur et al. 1995)

poly(styrene-b-dimethylsiloxane) (PS-b-PDMS) (Chu et al. 1995) that underscores the generality of coil-coil BCP phase behavior.

Liquid Crystals and Liquid Crystalline Polymers LCs are another class of materials that undergo self-assembly and form phase structures in the 1–10 nm length scale (de Gennes 1975; Kato et al. 2006; Tschierske 2007; Wang and Zhou 2004). However, unlike semicrystalline polymers, LC polymers (LCPs) possess excellent electro-optical properties that can be controlled using external stimuli (Collings and Hird 1997). Associating an LCP into a BCP results in LCBCPs that exhibit hierarchy in structure and also functionality. Prior to the discussion of LC BCPs, we shall first briefly discuss LCs and LCPs.

Liquid Crystals LC phase is a state of matter that is intermediate between crystalline solid and isotropic liquid states and is often referred to as mesomorphic (intermediate) state. LC phase is characterized by lack of long range positional and/or translational order in the system compared with crystals. Mesogens are generally anisotropic with aspect ratio greater than 4 (Wang and Zhou 2004). Generally, they consist of rigid molecules (such as aromatic rings) forming the core of the mesogen and the rings can be directly linked to one another or can be joined using linking groups such as –CH2CH2–, –CH
CH–, –CH=N–, etc. as shown in Fig. 3a (Collings and Hird 1997). The aromatic rings provide rigidity to mesogens and help maintain their anisotropic structure. Attached to the rigid aromatic core are aliphatic tails that

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a H3CO

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Spiralling polarization direction Smectic A

Chiral mesogen

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Front view

Side view

P Antiferroelectric

Back view

P Ferroelectric

Fig. 3 (a) Molecular structure of a typical calamitic liquid crystal mesogen. (b) Phase structures of most common LC phases. (c) Schematic representation of BCLCs in SmAP AF & F arrangement. (Adapted from de Gennes 1975; Collings and Hird 1997; Niori et al. 1996)

induce flexibility in the molecule and help modify the accessibility of the LC phases within moderate temperatures. In a mesogen, transformation from X to an LC phase can occur due to increase in temperature or due to a change in concentration. The former are called thermotropic LCs whereas latter are called lyotropic LCs. General LC mesogens can be either rod-like (calamitic LCs) or disc-like (discotic LCs) in shape. In an LC phase, the mesogens orient in a preferred direction called the director and is denoted by the unit vector, n^. The amount of orientational order in an LC phase can be defined using an order parameter, S. In a sample where all the molecules are perfectly oriented in a particular direction, S = 1 and for perfect disorientation, S = 0. For LCs, S varies between 0.3 and 0.8 (Collings and Hird 1997). In thermotropic LCs, S decreases with an increase in temperature due to the transition from X to I via LC phase. In different materials, this phase transition results in various unique correlations between the molecules that are identified as different LC phases (Collings and Hird 1997; de Gennes 1975; Zhou et al. 1989).

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Nematic (N) LC phase is characterized by the absence of long range translational order; however, all the molecules are parallel to the director. N phase is fluid-like in nature with high molecular mobility (Fig. 3b) (de Gennes 1975). Smectic (Sm) LC phase possesses higher degree of order compared to N. Sm phases are layered and possess definite interlayer spacing. Based on the molecular arrangement within these layers, different types of Sm phases are present. In Smectic A (SmA) LC phase, the mesogens within a layer do not possess positional order but their molecular axis is parallel to the layer normal. In Smectic C (SmC) phase, the molecules within layers are tilted with respect to the layer normal. SmC is optically biaxial and is considered as an important phase from application point of view in technology, because the molecules can be rotated about an axis using external field. Chiral SmC (SmC*) phase is a SmC phase with chiral mesogens and possess a helical twist from layer to layer. The polarizing direction varies along the twist whereas the tilt angle with respect to the normal remains constant. Other types of mesogens with nontraditional shapes have also been reported. For example, bent-core or banana LCs (BCLCs) are so named because there is a bend in the shape of the mesogen as shown in Fig. 3c. Niori et al. (1996) were the first to discover unique polar order formation in BCLCs and since then a variety of mesophases were discovered in these materials and were named as B1–B8 depending upon the sequence in which they were discovered. The advancements in the field of synthesis and characterization of phase structures of BCLCs was recently published in a review article by Reddy and Tschierske (Coleman et al. 2003; Heppke and Moro 1998; Link et al. 1997; Pelzl et al. 1999; Reddy and Tschierske 2005; Ros et al. 2005; Tschierske and Dantlgraber 2003). Like calamatic LCs, BCLCs also undergo self-assembly and form different LC phases but the uniqueness of BCLCs compared to linear mesogens is the additional complexity they impart to the LC phase due to the unique bent shape (Reddy and Tschierske 2005). For example, BCLCs possess a dipole moment in the direction perpendicular to the molecular axis and their unique shape facilitates their efficient packing. When these molecules, possessing a dipole moment, are packed into SmA layers, polarization of the layers can be achieved (denoted as SmAP where P stands for polar order) as shown in Fig. 3c. If the direction of polarization alternates between layers, it leads to anti-ferroelectric (AF) behavior whereas uniform polarization between layers leads to ferroelectric (F) behavior in the material. These two states can be switched by applying external field such as electric field. When the bent core molecules are tilted with respect to the Sm layer normal, SmCP (SmC denotes the mesogens are tilted and P denotes polar order) LC phases are obtained. Research in the field of SmCP LC phases has gained a lot of momentum in the recent years due to the wide variety of orientations and ordering achievable in these systems.

Liquid Crystalline Polymers Polymeric liquid crystals or liquid crystalline polymers (LCPs) are achieved by incorporating mesogens into polymer chains. Within an LCP, the rigid mesogen

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units may be present along the main chain (main-chain LCPs, MCLCPs) or as pendant groups attached to the side (side-chain LCPs, SCLCPs) using soft spacers such as methylene units as shown in Fig. 4a, b (Mao and Ober 1998; Walther and Finkelmann 1996). In SCLCPs, the mesogen can be attached to the backbone end-on or laterally. In MCLCPs, the presence of mesogens along the backbone imparts rigidity to the polymer chain whereas in SCLCPs the backbone is more flexible because the interactions with the mesogens are decoupled by the flexible spacers. This difference in the type of LCPs leads to differences in their properties, e.g., mainchain LCPs have higher modulus and transition temperatures compared to the sidechain LCPs. LCPs also exhibit mesophases similar to their small Mn counterparts although polymerization imparts better order and higher transition temperatures to LCPs. Although both Sm and N LC phases are observed in MCLCPs, Sm structures are more common in SCLCPs. The flexibility of polymer backbone facilitates packing of the mesogens into layers. Mesogen-jacketed liquid crystalline polymers (MJLCPs) are a unique class of SC-LCPs in which the mesogens are attached laterally along the polymer backbone without a spacer. Finkelmann et al. proposed that lateral attachment of mesogens along the polymer backbone restricts the rotational motion of the mesogen around its long axis and facilitates LC ordering (Hessel and Finkelmann 1986). By decoupling the mesogens and the backbone using spacers, LC ordering can be achieved in laterally attached LCPs although the backbone adopts a statistical chain conformation (Ban et al. 2014; Chen et al. 2010; Liang et al. 2010; Pan et al. 2007; Yu 2014; Zhou et al. 1987, 1988, 1989, 2010; Zhu et al. 2014). In MJ-LCPs, the size of the mesogen is larger than the backbone repeat unit. In such a situation, packing of the Fig. 4 Molecular architecture of (a) main chain, (b) side chain, and (c) mesogenjacketed liquid crystalline polymers

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mesogens is achieved by forcing the backbone to adopt an extended chain conformation due to steric hindrance of the mesogens. Thus macromolecular columns, with characteristic rigid rod-like structure, are formed where the polymer backbone forms the axis of the rod (Fig. 4c) (Zhou et al. 1987, 1988, 1989). Additionally, since the mesogens are attached to their “gravity centers,” the motions of the backbone were expected to have the least effect on the mesogen orientation and this decoupling was achieved without using spacers. N LC phase structures were observed by Zhou et al. both in monomers and polymers of 2,5-bis(p-methoxybenzoyl)oxy)styrene based MJ-LCP system (Zhou et al. 1987, 1988, 1989).

Liquid Crystalline Block Copolymers LCBCPs are formed by introducing LCPs into BCPs and since LC ordering occurs at 1-10 nm length scale which is an order of magnitude smaller than the dimensions of ordered nanostructures formed due to BCP microphase separation (10–100 nm), LCBCPs exhibit hierarchical structures with characteristic self-organization at multiple length scales. The overall phase behavior of LCBCPs is influenced by two competing interactions: BCP microphase separation and LC ordering. This results in unique phase structures (not attainable in coil-coil BCP systems) depending upon which of the two interactions dominates. In an LCBCP, the mesogens can be associated with the LC block of the BCP using either covalent or non-covalent interactions (Chen et al. 2010; Fischer and Poser 1996; Fischer et al. 1994; Mao and Ober 1998; Olsen and Segalman 2008; Walther and Finkelmann 1996; Yu 2014). Using MCLCPs or SCLCPs as one of the combining blocks in a BCP results in covalently bonded LCBCPs whereas noncovalently bonded LCBCPs consist of introducing the mesogen into the BCPs using weaker secondary interactions such as hydrogen bonding, ionic interactions etc. These two types of LCBCPs differ in the degree of rigidity of the LC block. Covalent bonding imparts more rigidity compared to the noncovalent interactions, and within the covalently bonded systems, MCLCPs possess more rigidity compared to SCLCPs. The term “rod-coil” (RC) is often used to describe these systems, and it emphasizes the difference in flexibility between the rigid LC block and the flexible coil block in a covalent LCBCP system. The increased rigidity contrast results in well-defined BCP nanodomains even in low Mn systems. Consequently, the LCBCP ODT occurs at a lower value of χN (higher temperature) compared to coil-coil systems. LCBCPs are a very promising system both from a technological and scientific point of view. Due to the different length scales of ordering, hierarchy in ordered structures can be obtained and by incorporating various groups within each of the blocks, different functionalities can be achieved at different length scales. From a scientific point of view, phase behavior of these systems will be different compared to a coil-coil BCP system because LCs are rigid and characterized by preferential orientation of the director. Novel self-assembling materials with complex phase structures are possible due to the competition between BCP microphase separation and LC ordering. In the following sessions, we shall discuss phase behavior and assembly behaviors of LCBCPs with covalent and noncovalent interactions.

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LCBCPs with Covalent Interactions: RCBCPs The thermodynamic phase behavior of coil-coil BCPs is characterized by the parameters f and χN. However, in RCBCPs, along with f and χN, additional parameters were required to account for the influence of the degree of asymmetry between the two blocks on the free energy and the resulting microphase separated structures (Matsen and Barrett 1998; Pryamitsyn and Ganesan 2004). Theoretical calculations have shown that one of these parameters is the parameter ν which is the ratio between the coil Rg and the rod length. The term ν accounts for the difference in scaling behavior of the two blocks because their aspect ratio varies differently with Mn. For a coil, an increase in the Mn increases its interfacial area (Acoil) at the intermaterials dividing surface (IMDS), whereas in a rod an increase in Mn increases its length while Arod remains constant. This induces instability at the interface and it must be taken into account in mapping RCBCP phase behavior. The second additional parameter is μN (Maier-Saupe interaction parameter), and it is a measure of the degree of orientational order that arises due to the rod-rod interactions. μ is the strength of the orientational interactions that favor the alignment of rods and it varies inversely with temperature (Pryamitsyn and Ganesan 2004). A number of research groups reported unique phase behavior in RCBCPs based on both theoretical calculations and experimental observations. Conventional BCP morphologies such as S, C, G, and L were observed in low Mn oligomeric rod systems (Lee and Cho 2001; Lee et al. 2000, 2001, 2004; Lee and Oh 1996; Lee and Yoo 2002). With an increase in Mn and frod, the rod-rod interactions became stronger resulting in the suppression of morphologies with curved interfaces (G, C, and S). Structures with flat interfaces became more predominant. Using SCFT, morphologies such as striped and hockey-puck shapes were calculated in the coil-rich samples whereas rod-rich samples exhibited zig-zag L and arrow-head structures (Pryamitsyn and Ganesan 2004). Within these unique morphologies, depending upon the orientation of the rod blocks, N, SmA, SmC and bilayered phases were theoretically calculated using free energy calculations in the SSL (Chen et al. 1995, 1996; Semenov 1985, 1986; Thomas et al. 1997; Williams and Fredrickson 1992). The WSL morphologies were predicted based on fcoil where samples with high fcoil transformed to microphase separated structures (from an isotropic phase) while at low fcoil, they transformed into N phase (Holyst and Schick 1992). Figure 5a, b shows the phase diagram of RCBCPs developed using SCFT approach by Pryamitsyn et al (Pryamitsyn and Ganesan 2004). Experimental observations corroborated theoretically calculated phase behavior. A variety of molecules that possess inherent rigidity were used as the rod block to synthesize RCBCPs. Both polymers and oligomers were used as the rod block. The rod block, of most of the RCBCP systems reported, consists of molecules that are either helical rods, mesogenic rods or conjugated rods. Synthetic polypeptides adopt conformations that form α-helices and β-sheets that are inherently rigid in nature. A unique double hexagonal structure (Fig. 5c) was observed by Klok et al. who investigated the phase behavior of oligopeptide based diblock oligomers of oligostyrene (St) (with degree of polymerization (DP) of 10)

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Fig. 5 Theoretical phase diagram of RCBCPs (f represents volume fraction of the coil) with (a) ν = 0.15 and (b) ν = 0.25. (c) Double hexagonal hierarchical structure and (d) hockey-puck shaped aggregates. (Adapted from Pryamitsyn and Ganesan 2004; Klok et al. 2000; Schleuss et al. 2006)

and oligo (γ-benzyl-L-glutamate) (OBLG) (DP = 10 and 20) (Klok et al. 2000; Klok and Lecommandoux 2001). OBLG predominantly possesses α-helical conformation in the sample with 20 repeating units. The RCBCPs form hexagonally close packed C spaced 4.3 nm apart. Within these larger C, the helical rods are arranged forming a hexagonal close packed C lattice with 1.6 nm spacing. The smaller peptide C were oriented with their C long axis perpendicular to the axis of the larger C. Increasing the asymmetry by increasing fcoil leads to the formation of unique hockey-puck shaped aggregates as observed in an oligomeric hepta( p-benzamide)-b-poly(ethylene glycol) system (DProd = 6 and DPcoil = 110 and 45) by Schleuss et al. (Fig. 5d) (Schleuss et al. 2006). These hockey-puck shaped aggregates were enclosed in spherical micelles of diameter 35 nm and were similar to the theoretical puck shaped aggregates proposed by Williams et al. and Ganesan et al. (Pryamitsyn and Ganesan 2004; Williams and Fredrickson 1992). The core of the spherical micelles consists of puck shaped aggregates (10 nm  4.4 nm  2 nm) surrounded by PEG corona. Based on the results obtained from scanning probe microscopy (SPM) and dynamic

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light scattering, the authors proposed that the core contains bilayer molecules that are stabilized by intramolecular hydrogen bonding and π-π interactions. The uniformity in diameter of the PEG corona, although the core is anisotropic, implies that PEG coils near the core stretch in order to fill the space and the ones away from the core are more coil like. Shorter coils (DPcoil = 45) could not form the shell around the aggregates and resulted in the formation of rod-like micelles. Using a variety of hydrophobic and hydrophilic polypeptide chemistries ((poly (γ-benzyl-L-glutamate) (PBLG), poly(γ- L-glutamic acid) (PLGA), poly(Z-LLysine) (PZLL), poly(N-trifluoroacetyl-L-Lysine) (PNTLL), and poly (hydroxyethyl-L-glytamine) (PHLG)) as rod blocks and PS, PB, and polysarcosine as coil block, Gallot et al. investigated the phase behavior of these RCBCPs both in solution and melt (Douy and Gallot 1982; Gallot 1996). Seven different phase structures were reported in these systems based on the orientation of the rods although Sm and N phases were predominant. All the BCP systems demonstrated the remarkable preference of the L phases in RCBCPs. In L forming RCBCP sample, the α-helical rods were arranged perpendicular to the IMDS and possess hexagonal in-plane symmetry. Based on layer thickness and extended chain length of the helices, it was observed that the helices generally adopted a folded chain conformation (Fig. 6a) (Douy and Gallot 1982). Folded conformation imparts more interfacial area for the coil block and is hence preferred. In general, interdigitation, tilting, and folding of the rods provide more interfacial area to the rod blocks thereby reducing the coil stretching penalty. Hexagonal-in-lamellar morphologies were also observed in rod-coil-rod triblock samples where PB and polysarcosine formed the coil block, and PBLG, PNTLL, and PHLG formed the rod block. However, in these samples, within the Sm layers the helical rods were tilted at an angle with respect to the layer normal and the tilt angle increased as fcoil increased (Fig. 6b). These tilted rods were, however, still arranged in a hexagonal lattice with the same lattice constant although the thickness of the rod layer decreased as the tilt angle increased. Theoretical calculations of Semenov et al. showed predominantly N, SmA, and SmC structures in their model RCBCP system (Semenov 1986). Halperin’s theoretical prediction showed tilting of the Sm layers with an increase in the fcoil in order to offer more interfacial area to the coil to reduce the coil stretching penalty and is in agreement with the experimental observation of Gallot et al (Halperin 1989, 1990). While the influence of fcoil on the overall morphology was investigated by Gallot and coworkers, Losik et al. reported the influence of solvent (dimethyl formamide, DMF) on the conformation of the polypeptide (Losik et al. 2004). In polypeptidebased RCBCPs with PBLG and PZLL as the rod block, with constant fcoil (PS DP = 52) and DP PBLG = 104 and PZLL = 111, they observed hexagonalin-L hierarchical structures. Within the rod blocks, PZLL adopted a fully stretched helical conformation whereas PBLG formed helices that were twice folded. This folding effect was attributed to the ability of the solvent (DMF) to penetrate the PBLG and soften the helical backbone by reducing the hydrogen bond interactions. This also led to a decrease in the order of the hexagonal lattice in PBLG rods. PDI of the system also influences the BCP structures. Schlaad et al. also investigated the effect of variation in PDI on the interface using PZZL based RCBCPs (Schlaad et al. 2004).

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a

b

O CH2 CH

n

C N

m

(CH2)5CH3

c

d ∧

n ∧

p

∧

∧

∧

n

p

n ∧

p

D

D

200 nm

400 nm

Fig. 6 Schematic representation of the phase structure of poly(vinyl-b-peptide) BCPs: (a) folded chain (Douy and Gallot 1982) and (b) tilted layered morphology. (c) Zig-zag lamellae and (d) arrow head morphologies of PS-b-PHIC. (Adapted from Chen et al. 1996)

In this system, hexagonally packed polypeptide α-helices were observed in zig-zag L morphology where the main axis of the helix is oriented perpendicular to the BCP interface. The zig-zag nature of the BCP L was due to the kinks that are formed due to the fractionation of helical rods according to their length. Samples with large PDI had relatively planar interface between the kinks and were characterized by less number of kinks per unit volume whereas samples with narrow and moderate PDI (1.01–1.27) exhibited a higher number of kinks. Although the polypeptide systems form good rod blocks, they exhibit different configurations of the molecules (α-helices, β-sheets) and are also observed to undergo folding, which makes interpretation of their phase behavior difficult. Using poly(hexyl isocyanate) (PHIC), which possesses only stiff helical rod (persistence length ~ 50 nm–60 nm) conformation, as the rod block and PS as the coil very unique phase structures were reported by Ober and Thomas groups (Chen et al. 1995, 1996; Thomas et al. 1997). A series of samples of PS-b-PHIC RCBCP system with Mw varying from 70,000–1,800,000 g/mol and frod = 0.42–0.98 were investigated and unique morphologies such as wavy L (lenticular aggregates), zig-zag L, and arrow-head shaped domains were observed. A sample with fPHIC = 0.42 showed lenticular aggregate structures and at fPHIC = 0.73–0.90 the samples displayed alternating layers of PS and PHIC arranged in a unique zig-zag manner (Fig. 6c) (Chen et al. 1995; Chen et al. 1996; Thomas et al. 1997). Within the PHIC domains, the rods were arranged forming a long range Sm order. Based on the Mn and d-spacing calculations, they concluded that the Sm layers are interdigitated. The PHIC rods are tilted with respect to the interface thus forming SmC LC phase and the angle of tilt increased with an increase in fcoil. At very high concentrations of the rods ( frod = 0.98), the system exhibited an arrow-head shaped PS morphology with a flip in the orientation of the head by 180 in every alternative layer (Fig. 6d) (Chen et al. 1995, 1996; Thomas et al. 1997). At high frod, the domain spacing of the coil was smaller than the Rg of PS, leading to the formation of inhomogeneous Sm monolayer phase termed as SmO (similar to the small molecule LCs where the orientation of director flips between layers). In a RCBCP of poly(styrene-b-

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(3-(triethoxysilyl)propyl isocyanate)) (PS-b-PIC) (DPPS ~ 1900 and DPPIC ~ 100) and with fPS = 0.9, Park et al. observed anisotropic nanoscale objects in the shape of parallelograms that form Sm ordering in solution cast films (Park and Thomas 2004). Using theoretical calculations, it was shown that these parallelograms, with a side of at least 45 nm, had a tilted bilayer arrangement of the PIC rods. The authors propose that interdigitation of the rods is also possible. Within these nano-objects, the PIC rods are oriented uniaxially and these nano-objects have a 25–45 angle between the sides. Rods made of mesogenic molecules also possess inherent rigidity as a result of their molecular chemistry. A lot of interesting morphologies were observed in RCBCPs containing mesogenic rods although most of the mesogens were based on low Mn molecules. Stupp’s group investigated the phase structures in mesogenic oligomers (formed by attaching an azo dye to a rigid biphenyl carboxylic acid) of length 6 nm attached to flexible polyisoprene (Mw = 3200 g/mol). Good control over the polydispersity of the rod molecule allowed the researchers to explore the subtle competition between the microphase separation and LC ordering processes by altering the chemistry and Mn of the molecules (Radzilowski et al. 1993, 1997; Radzilowski and Stupp 1994; Stupp 1998; Stupp et al. 1997). At frod = 0.36, the rigid mesogenic rods formed stripes arranged in layers inside the PI matrix. At similar f values, hexagonally close packed C morphologies are observed in coil-coil systems. In the present RCBCP case, although the stripes possess a hexagonal symmetry, the nature of the orientation of the rods within the stripes was not clear (Fig. 7a) (Lee et al. 2001). Qualitatively, these stripe-like patterns are similar to the broken L phases as reported in the numerical calculations by Ganesan et al. (Pryamitsyn and Ganesan 2004). As the frod decreased to 0.25, the morphology transformed from stripes into super lattice aggregates (7 nm in diameter) that are oriented parallel to the substrate (Fig. 7b)(Lee et al. 2001). These aggregates are arranged with hexagonal in-plane symmetry in a PI matrix. Theoretical calculations by Fredrickson et al. and Ganesan et al. also showed similar puck-like morphologies (Pryamitsyn and Ganesan 2004; Williams and Fredrickson 1992). The authors claim that the observation of a hexagonally packed lattice at such highly asymmetric f suggests the nonuniformity in distribution of the coronal coils around the core. The coils grafted to the edge of the core adopt a radial conformation whereas the ones at the center are more elongated. The authors suggest that this resulted in nonuniformity in space filling that manifests in the distortion of an otherwise BCC lattice. In a triblock RC architecture with oligostyrene and oligoisoprene coils and three biphenyl units as rods, Radzilowski et al. reported unique mushroom-shaped nanostructures where an aggregated packet of rods form the stem and the oligomeric coils splay forming the head of the mushroom. About 100 rod molecules aggregate into a mushroom (the size is limited by the repulsive interactions of the oligomers coils that form the head of the mushroom) (Fig. 7c) (Stupp et al. 1997). These mushrooms further stack to form a layered structure. The rods orient normal to the substrate leading to self-organized films with polar and nonpolar bottom and top surfaces. Keeping the length of the rod constant, the authors observed a disruption of the layered structures with an increase in the fcoil due to the steric repulsions between

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CL Et

N

O2N

NCH2CH2OOC

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COO

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OOC(CH2)7COO

COO

OOC

OOC n

b

[

[[

CH2 CH

– 9

n

[

COO – m n+m=9

OOC

OOC

OH

c

Fig. 7 Schematic representation of (a) stripes and (b) super-lattice aggregates of low Mn mesogenic units. (c) Mushroom-shaped aggregates. (Adapted from Lee et al. 2001; Stupp et al. 1997)

the bulky coils, whereas keeping the fcoil constant and increasing the length of the rods resulted in stabilization of mushrooms. Introducing a bulky dendritic wedge at the end of the rods increased the steric hindrance of the packing of the mushrooms and the morphology was completely altered resulting in ribbon-shaped structures of width 10 nm and thickness 2 nm. Lee et al. systematically varied the chemical structure of the rod and coil blocks and reported very interesting phase behavior in a rod-coil oligomeric system where the rod units are made of mesogens (Jin et al. 2004; Lee and Cho 2001; Lee et al. 2000, 2001, 2004; Lee and Yoo 2002). LC ordering was predominant in longer rods (two biphenyl units) compared to smaller rods (single biphenyl unit) due to the increased lateral interactions between the rods (keeping the coil segment constant) and leading to the formation of L structures. The rod domains transform into hexagonal columns as fcoil increased. This transformation can also be brought

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about by keeping the fcoil constant and changing the coil chemistry from PEO to PPO due to more space requirements of the bulkier PPO coil. In a system with PPO as the coil block and the rod made of two biphenyl units with constant length of the rod block, the authors observed SmA and SmC structures of the rods at symmetric fcoil, forming layered structures. An isotropic bicontinuous cubic phase and hexagonally closed packed C of the rods were observed with an increase in the asymmetry ( fcoil) (Fig. 8a, b) (Lee et al. 2001). Within the minority domains, the rods formed aggregates with square cross-section. Further increase in fcoil led to the formation of discrete spherical micellar structures. In all these phases, the rods in the minority blocks possess a bilayered Sm structure. A slight increase in the rod length by using an extra phenyl linkage showed remarkably different phase structures stabilized by improved LC ordering. Unique honey-comb like layered structures, where PPO formed perforations of 6.5 nm arranged 10 nm apart with hexagonal in-plane symmetry (Fig. 8c) (Lee et al. 2001). Furthermore, attaching PPO coils on either end of the rod resulted in the formation of oblate-shaped aggregates arranged in a body-centered tetragonal (T) lattice (Fig. 8d) (Lee et al. 2001). Based on X-ray scattering data, the authors calculated 84 molecules per aggregate in the T packed structure (diameter 5 nm and length 3 nm). Based on theoretical theoretical calculations, the authors propose that the oblate shape of the aggregates is responsible for the unique T arrangement as this structure imparts maximum packing density to the system. Further increasing the length of the rod by introducing a phenyl group into the molecule resulted in a hexagonally ordered mesophase that transformed into a T micellar phase only at high temperatures. Further increasing the length of the rod completely suppressed the formation of the T phase and resulted in L morphologies with in-plane hexagonal symmetry in the rod block. The strong effect of coil stretching penalty was demonstrated by increasing the grafting density of polymer chains per interface. This was achieved by serial combination of (rod-coil)n units with varying n values. In bulk, when n = 1, L crystalline structures were observed and when n = 2 and n = 3, the morphology changed to 2D rectangular crystalline and hexagonal columnar structures, respectively. As n increased, the grafting density per interface increased and this resulted in the transformation of sheet like rod domains into individual aggregates arranged in a 2D lattices so as to reduce the stretching penalty of the coils. Rod blocks made up of conjugated molecules have also been used to form RCBCPs. Conjugated molecules consist of rigid cores that impart LC character to the molecule. Aggregation of these molecules to form LC phases helps in orientation of the delocalized π electrons making them ideal candidates for photonic and electronic applications. Organizing these materials into ordered macroscopic structures is necessary to exploit their properties in the bulk. RCBCPs formed by π conjugated oligomers and polymers (like oligo- or poly- p-phenylenevinylene (PV), ( p-phenylene) (PP), thiophenes, phenyl quinolines (PQ), phenylene ethynylene (PhE)) thus led to the formation of materials with structural and functional hierarchy. PPP and PPV have potential applications as components in light emitting diodes (LEDs) or solid state light emitting cells (SECs), photovoltaic cells etc. (Hide et al. 1996; Sirringhaus et al. 1998). Using these polymers as rod blocks in

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OOC

CH3CH2OOC

a

O

O H n

n = 10-15 b

n = 15-20

c CH3(CH2)2O

OOC

COO

O

O CH3 n

d CH3 O

O n

OCH2

CH2O

O

O CH3 n

82.7 Å 30 Å

89.9 Å 50 Å

Fig. 8 Schematic representation of (a) bicontinuous cubic phase, (b) hexagonally closed packed C, (c) layered honeycomb, and (d) oblate structures arranged in body-centered tetragonal lattice observed in rod-coil oligomeric system by Lee et al. (Adapted from Lee et al. 2001)

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an RCBCP led to improvement of their properties compared to the homopolymers due to enhanced ordering (Liang et al. 2007). Using low Mw oligo PV as the rod block and PI as the coil block with varying wrod (weight fraction of rod, 0.12–0.4), Li et al. reported L morphologies with bilayered supramolecular structures within the OPV block (Li et al. 1999). Introducing amphiphilic nature to these BCPs by using PEG instead of PI, Wang et al. observed long cylindrical micelles (15–18 nm core diameter) with OPV core formed in water/ THF solvents and in thin films (Wang et al. 2000). Further changing the coil to hydrophobic PPG led to the formation of well-oriented fiber patterns on mica substrate (Wang et al. 2004). The orientation of these fibers was attributed to the interactions between the coils and the substrate where longer coils led to better oriented fibers. In PS-b-PPQ system, Jenekhe et al. reported unique supramolecular phase structures such as hollow S, L, hollow C, vesicles, etc. depending upon the type of solvent used and the rate of solvent evaporation (Fig. 9a) (Chen and Jenekhe 2000; Jenekhe and Chen 1998, 1999). By changing the type of solvent, they were able to obtained hollow hard spheres (rod core) and hollow soft spheres (coil core). At a macrolevel, the hard spheres undergo self organization to form 2D hexagonally ordered arrays of spherical holes formed due to the condensation of the solvent during annealing. The nano hollow spheres were also used to encapsulate molecules such as fullerenes (Chen and Jenekhe 1999). Sary et al. explored the relationship between χ and μ in P4VP-b-PV RCBCP system (Sary et al. 2007a, b). The ratio μ/χ determined whether the LC ordering or microphase separation dominated the overall morphology. In their system (χN > 15–20 and μ/χ < 4), Sary et al. observed the domination of the microphase separation and the PPV rods were confined into L, C, and S domains. However, the C and S domains were highly distorted and exhibited poor long range order. Within the L, annealed PPV domains showed SmC2 (bilayered SmC) phase with a tilt angle of ~52 . This SmC2-in-L transformed into L which further transformed into isotropic phase with increase in temperature. In the PPV hexagonal C structure, the molecules were arranged radially, perpendicular to the C axis. In a PSb-PPV RCBCP with frod = 0.17 which forms homogeneous isotropic structures, blending with PPV homopolymers resulted in the formation of SmC2 structures due to the π–π conjugation between the homopolymer PPV and the rod block of the RCBCP (Fig. 9b) (Sary et al. 2007a). These rods were tilted at an angle of 54 with respect to the layer normal. Blending of up to 50 wt.% PPV did not result in any macrophase separation. They attribute the increase in the Sm ordering upon blending to the increased rod-rod interactions. By attaching a cyano substituted PV chromophore segment to the rod segment of their mushroom forming oligomeric system, Stuppet al. observed the fluorescent properties of the molecules (Pralle et al. 2000). The observation of strong polar ordering in layered mushroom structures of ~5 nm diameter after conjugating indicates the high thermodynamic stability of the mushroom structures. Strong photoluminescence and piezoelectric behavior and mechanical adhesion were reported in these conjugated molecules.

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Fig. 9 Schematic representation of (a) hollow sphere structures formed by PS-b-PPQ and (b) SmC2-in-L hierarchical structure in PS-b-PPV system. Phase diagram developed from experimental observations of PPV-b-PI system in the (c) weakly segregated and (d) moderately segregated samples. (Adapted from Jenekhe and Chen 1998; Olsen and Segalman, 2005, 2006)

Segalman et al. studied the room temperature and high temperature phase behavior of weakly segregated PI-b-PPV (alkoxyphenylenevinylene) RCBCP system consisting of PPV rods and PI coils with varying fcoil (0.42–89) but constant PPV rod length (Mn PPV = 3545 g/mol) (Olsen and Segalman 2005, 2006). All the samples exhibited L phases at low temperatures. However, with an

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increase in temperature, all the samples (except the sample with highest fcoil which is isotropic throughout the temperature range) transformed into N phase that further transformed into an isotropic phase. Using a combination of X-ray and light scattering, optical and electron microscopy results, the authors developed an experimental phase diagram for the weakly segregated system, which is shown in Fig. 9c (Olsen and Segalman 2005, 2006). The width of the N phase region in the phase diagram increases with increase in the frod. Within the PPV domains, the rods could be arranged either in monolayers or tilted bilayers. However, they did not observe either the hockey puck morphology at high fcoil or the zig-zag structures at high frod. The authors claim that a higher aspect ratio is required for the formation of puck-shaped morphologies. Accordingly, greater aspect ratio leads to greater elastic energy of bending which in turn induces greater coil stretching and the puck formation is a way of reducing this chain stretching. The high temperature phase behavior of these systems follows the trend: low fcoil L ➔ N ➔ Isotropic and high fcoil L ➔ Isotropic. The same research group also reported the phase behavior of PPV-b-PI system with higher Mw PPV (5600 g/mol) with increased geometrical asymmetry between the rod and coil blocks (Olsen and Segalman 2007). The increase in Mw resulted in moderately segregated BCP system compared to their previous weakly segregated system (Mn PPV = 3545 g/mol). Keeping the PPV length constant, the authors reported the phase structures of a series of sample with varying fcoil (0.31–0.91). At moderate fcoil, the authors observed L phases, however, increasing the fcoil to 0.81 led to the formation of PPV nanodomains packed in a hexagonal lattice. PPV rods inside these domains do not possess any preferred orientation with respect to the interface. These nanodomains had rectangular interfaces (~32 nm wide and ~12 nm thick) corresponding to a packing of ~1000 rods per cluster. The packing of the rods improved as the frod increased. The sample with fcoil = 0.81 is considered as the boundary between the transition from L to hexagonally close packed aggregate structure (H). Faster quenching from melt resulted in the retention of these H aggregates, which were stable at low temperatures, whereas slower annealing below the ODT resulted in OOT into L phase. Since entropic factors of coil stretching prefer H phase compared to L phase, the OOT into L phases suggested that in these samples, the rod orientation dominates the BCP microphase separation leading to L phase formation. The H phase is present only at higher fcoil. The authors proposed that within the L phases, the rods are oriented perpendicular to the L interface. Heating the samples resulted in N phase before isotropization in all the samples, irrespective of the fcoil. This is different from their observation of absence of N phase at high fcoil in weakly segregated samples and is attributed to the increase in the rod-rod interactions due to higher Mw. Based on their experimental observations, Segalman et al. developed a phase diagram for the PPV-b-PI system in the moderate segregation limit as shown in Fig. 9d. L phases dominated the phase diagram at low and moderate fcoil, and at high fcoil, the H phase was observed. Increase in PPV Mw led to the formation of N phase in all the samples irrespective of the fcoil.

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LCBCPs with Covalent Interactions: Mesogen Jacketed Liquid Crystalline Block Copolymers As previously discussed, by laterally linking the “waist” of LC mesogens directly to polymer backbones (without spacers), MJLCPs can be achieved (Gopalan and Ober 2001; Pragliola et al. 1999; Wan et al. 1998; Zhang et al. 1999; Zhou et al. 1989). In most of the MJLCP systems, rigid columns of the mesogens are formed due to the strong interaction between the side chain mesogens and the polymer backbone. These rigid columns pack together forming columnar nematic (ΦN)/hexagonal (ΦH) phase, and it is these supramolecular columns, rather than the individual mesogens, that possess the orientational order. The rigid nature enables the MJLCPs to serve as rods and form a new type of rod-coil block copolymers. Morphology and rheological behavior of rod-coil poly(styrene)-block-poly(2,5-bis-(4-butyl-benzoyl) oxystyrene) (PS-b-PBBOS) has been reported by Ober et al. (Gopalan et al. 2003a, b). A series of MJLCP-based rod-coil block copolymers, poly(styrene-block-(2,5-bis [4-methoxyphenyl]oxycarbonyl)styrene) (PS-b-PMPCS) have been synthesized and their solution self-assembly behavior has been investigated (Tu et al. 2000; Wang et al. 2005; Yi et al. 2004). Symmetric PS-b-PMPCS with relatively low Mw forms simple lamellar phase with the supramolecular PMPCS rods aligning parallel to the lamellar normal. Each LC layer consists of approximately two layers of PMPCS and a bilayer SmA phase was thus proposed (Li et al. 2004). The phase structure of asymmetric PS-b-PMPCS BCP is much more complicated (Tenneti et al. 2005). For example, PS171-b-PMPCS32, with fPMPCS ~ 0.37, a perforated layer (PL) structure was observed. Figure 10 shows the TEM micrographs of the samples that were microtomed normal to y, x, and z directions defined in Fig. 11 (Tenneti et al. 2005). From these images, it is interesting to observe that the dark areas are not continuous layers. This discontinuity of the dark layers in both xz and yz planes suggests the formation of a PL structure. Due to the large volume fraction of PS ( fPS~0.63), PS punctuates the PMPCS layers and forms the perforation. The in-plane symmetry of the PL phase can be unambiguously determined by viewing thin section of the sample along z direction and it has been shown that a tetragonal perforated layer (TPL) structure is formed in PS171-b-PMPCS32 rod-coil BCP sample, although most of the reported PL structures in BCPs have hexagonal symmetry (HPL) (Forster et al. 1994; Fredrickson 1991; Hamley and Bates 1994; Laradji et al. 1997; Qi and Wang 1997; Zhu et al. 2001). Based on the TPL model, PS molecules punctuate and form isolated “islands” in the PMPCS layer, and these in-plane isolated domains obey a square lattice symmetry with a = b = 29.6 nm. Therefore, the lattice of the perforation can be estimated to be a = b = 29.6 nm and c = 42.7 nm. Figure 11 shows the schematic representation of the hierarchical structure of the TPL structure of PS171-b-PMPCS32: strong interaction between PMPCS mesogens and the backbone leads to a near extended conformation of the PMPCS, which form the rods. This rod-coil system self assembles into the TPL structure where PS punctuates the PMPCS layers. The PMPCS rods are parallel to the layer normal of the BCP structures.

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Fig. 10 TEM micrographs of thin sections of PS171-b-PMPCS32. The sample was sectioned along the xz (a), yz (b), and xy (c) planes followed by RuO4 staining for 30 min. Insets show their FFT pattern (Tenneti et al. 2005)

Fig. 11 Schematic representations of asymmetric PS-b-PMPCS TPL hierarchical structure. Note the rods are parallel to the lamellar normal (Tenneti et al. 2005)

Another type of MJLCP BCPs has been syntheisized using bent core LCs (BCLCs) (Chen et al. 2006; Tenneti et al. 2008). Poly[styrene-block-{3,5-bis [(40 -((400 -tetradecanoylbenzoyl)oxy)benzoyl)oxy]styrene}] (PS-b-PTBOS) as the rod (Chen et al. 2006). The five-ring mesogen ensures a rigid core of the rod with a relatively large diameter while the 14-C tails of the mesogen render a relatively thick “shell.”

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Fig. 12 Three different hierarchical structures (a) ΦH-in-L in PS202-b-PTBOS35, (b) ΦN-in-PL in PS567-b-PTBOS24, and (c) ΦN-in-S in PS160-b-PTBOS126. LC symmetry breaking in b and c is due to BCP self-assembly (Chen et al. 2006)

For symmetric PS202-b-PTBOS35, fPTBOS ~ 0.58), a columnar-hexagonal-inlamellar (ΦH-in-L) hierarchical structure was formed (Fig. 12a) (Chen et al. 2006). Interestingly, competition between liquid crystallinity and BCP self assembly was observed in two types of asymmetric PS-b-PTBOS samples. PS567-b-PTBOS24 possesses an fPTBOS ~0.25 and it represents a PS-rich RCBCP. A hierarchical structure of columnar-nematic-in-perforated layer (ΦN-in-PL) was observed (Fig. 12). This LC symmetry breaking is possibly because in PL structures, the coil chains stretch and perforate the rod layer, there thus exists lateral repulsion of the coil chains as shown in Fig. 13 (Chen et al. 2006). Because the rods and the coils are covalently linked, this lateral repulsion of the coil chains further renders a splaying stress field on the LC rods, which forces the rod to bend thereby reducing the LC order. In order to confirm this, two blend samples were prepared and the PTBOS volume fractions were controlled to be 20% and 14%, respectively (PS567-bPTBOS24–20 and PS567-b-PTBOS24–14). As shown in Fig. 13b, compared to the WAXD pattern of pure BCP PS567-b-PTBOS24, the diffraction peaks of the blends became much sharper and the higher order reflections can be clearly seen. The LC order was thus dramatically increased upon blending, confirming that the ΦH LC symmetry was restored. The soft shell of the rods also dramatically influences the assembled structure in PTBOS-rich BCPs. In PS160-b-PTBOS126 ( fPTBOS = 0.86) film, PS forms spherical domains in the PTBOS matrix while the LC order was reduced and instead of ΦH, ΦN phase was formed. The hierarchical structure is ΦNin-S (Fig. 12). This clearly indicates that the curved IMDS reduces the order of LC and the LC symmetry breaking is because of the incompatibility of the translational symmetry of LC and the curved IMDS.

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Fig. 13 (a) Schematic representation of the stress releasing process in the ΦNin-PL hierarchical structure before (i) and after (ii) blending with low molecular weight PS. (b) WAXD patterns show that the LC ordering was dramatically enhanced in the blend samples with 20% and 14% PTBOS (1, 2) compared to the pure BCP (3) (Chen et al. 2006)

LCBCPs with Covalent Interactions: Side-Chain Liquid Crystalline Block Copolymers (SC-LCBCPs) SC-LCBCPs can be formed by introducing SCLCPs as one of the blocks in a BCP. Decoupling the strong influence of the mesogens on the polymer backbone showed a remarkable influence on their phase behavior. Compared with RCBCPs, typically SC-LCBCPs exhibit remarkable improvement in the LC ordering due to the spacers that decouple the backbone and mesogen conformations. This decoupling effect not only reduces the “rod-like” effect of the LC block but also contributes to a better understanding of the mutual interplay between LC ordering and microphase separation processes. Furthermore in SC-LCBCPs, since the mesogen possesses more flexibility, it can adopt a favorable orientation with respect to the IMDS (Fischer and

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a

b

Long spacer (Homeotropic)

c

Short spacer (Homogeneous)

Coll Cylinders

30

Lamellar

LC-Cylinders

25

c1N

20

(c1N)C

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(c1N)1–N

10 5 0.1

0.2

0.3

0.4

0.5

fA

0.6

0.7

0.8

0.9

Fig. 14 (a) Homeotropic and (b) homogeneous anchoring of the mesogens with the IMDS in SC-LCPs. (c) Theoretical phase diagram of SC-LCBCPs developed by Shah et al. (Osuji et al. 2000; Shah et al. 2008)

Poser 1996; Fischer et al. 1994; Gallot 1996; Walther and Finkelmann 1996). Homogeneous or homeotropic anchoring is possible depending upon whether the mesogen is oriented parallel or perpendicular to the IMDS, respectively (Osuji et al. 2000). The influence of the length of the spacers in determining the homogeneous or homeotropic anchoring of mesogens with respect to the BCP interface has also been studied (Fig. 14a, b) (Osuji et al. 2000). These orientations of the mesogens in turn influence the orientation of the Sm layers. A variety of SC-LCBCP systems were investigated with cholesteryl-, biphenyl benzoate core-, biphenyl ester core-, fluorinated-, and azobenzene-based LC mesogens side-attached to butadiene-, methacrylate-, isoprene-, and siloxane-based polymers and their phase structures have been reported (Al-Hussein et al. 2005; Ansari et al. 2003; Anthamatten and Hammond 1999; Anthamatten et al. 1999; Fischer and Poser 1996; Fischer et al. 1994; Gallot 1996; Hamley et al. 2004, 2005; Mao et al. 1997; Osuji et al. 1999, 2000; Poser et al. 1998; Tenneti et al. 2007; Tokita et al. 2007; Verploegen et al. 2007; Walther and Finkelmann 1996). Conventional BCP structures such as L, C, and S were observed in these samples, and the LC phase was predominantly Sm (both SmA and SmC) in the case of L and C forming systems whereas, N LC phases were observed at highly asymmetric fcoil. Shah et al. presented

7

Structure and Assembly of Liquid Crystalline Block Copolymers

199

theoretical analysis of the phase behavior in SC-LCBCPs as a function of Mn, rod length, and f of the blocks using SCFT and is shown in Fig. 14c (Note: N in the figure represents degree of polymerization and not N LC phase) (Shah et al. 2008). They observed that the segregation strength for microphase separation is much weaker than that for I to N (nematic) LC transition (in the figure, (χ1N)C < (χ1N)I-N) indicating that microphase separated structures are necessary to obtain ordered LC phases. C of coil in LC matrix was observed in samples with low fcoil in which the mesogen axis was parallel to the C axis. This kind of orientation was facilitated by the stretching of the polymer backbone away from the interface. Alternative arrangement of mesogens, oriented tangential to the C, would lead to defects in the structure. L phases were observed at symmetric f values where the mesogens were aligned parallel to the IMDS. Compared to the coil-coil phase diagram, theoretical SC-LCBCP phase diagram is asymmetric. At higher fcoil, C of the LC phase, with the mesogens aligned parallel to the C axis was the most favored morphology. In a PS-b-PI system where PI is functionalized with an azobenzene-based mesogen, Ober and Thomas et al. reported the phase behavior as a function of fPI-LC with morphologies including L, PS C in PI-LC matrix and PI-LC cylinders in PS matrix (Fig. 15) (Osuji et al. 1999, 2000). The LC block showed a TSmA-I at 171  C. In L samples and C forming samples where PS is the majority phase, the authors observed PI-LC forming SmA type of LC phase with homogeneously anchored mesogens. In shear-oriented C forming samples where PI-LC formed the majority phase, the authors observe a transformation of the orientation of the cylinders as a function of shearing temperature. When sheared in the Sm LC temperature, the cylinders were oriented with their long axis parallel to the neutral direction (perpendicular to the shear direction and in the plane of shear, also known as transverse orientation) and the Sm layers were oriented with their layer normal parallel to the neutral axis. The cylinders reoriented to adopt a parallel orientation (preferred orientation for coil-coil systems) when sheared in the TI. The authors propose four different possible models that are possible in a C forming LCBCP system: (where the LC forms the matrix) parallel-transverse, perpendicular-parallel, parallel-parallel, and transverse-perpendicular (Fig. 15). In this notation, the first word indicates the orientation of C axis and the second word is the orientation of the Sm layers. Among these possible orientations, the transverse-perpendicular orientation is the most favorable morphology compromise where both the Sm layers and the cylinders are not under strain and this morphology also satisfies the homogeneous anchoring condition of the LCs. Anthamatten et al. studied the room temperature and high temperature phase behavior of PS-b-PMMA based SC-LCBCPs where the methacrylate contained a chiral biphenyl benzoate mesogen attached using six and ten carbon alkyl spacers (Anthamatten 2001; Anthamatten and Hammond 1999; Anthamatten et al. 1999, 2001). The phase diagram (shown in Fig. 16) for this system showed predominantly L phase due to the layered nature of the Sm LC phase (Anthamatten et al. 1999). At lower fLC ( tension

F Upper fixture Ω -> Ω’-> Ω”

Fig. 4 (a) Scheme of the experimental setup. The sample is placed between two surfaces. The shear strain is transmitted to the sample via molecular contacts with the lower surface which is animated by an oscillatory motion of given frequency ω and amplitude γ 0. The shear stress is communicated along the sample thickness and is transferred via molecular contacts to the second surface coupled to a force (here a torque) sensor (real imposed strain geometry). (b) Simplified scheme of the transmission chain of the information in dynamic relaxation (in imposed strain geometry): the transmission of the shear strain to the sample and the transmission of the shear stress strain of the sample to the sensor are entirely tributary of the interaction forces between the liquid and the surface onto which it is deposited. This transmission chain can be formalized as follows: Ω is the imposed shear torque. Ω0 is the shear torque transmitted to the sample. Ω0 = Ω
losses from the surface to the sample (slip). Ω0 is the shear torque received at the sensor. Ω00 = F(Ω0 )
losses of the sample at the surface where F is the transfer function by the sample.

where the shear elastic modulus G0 (ω) and the viscous modulus G00 (ω) are the inphase and out of phase components, respectively. This stress response is interpreted as the Fourier Ð 1 transform of the stress relaxation Ð1 function G(t) of the quiescent state: G0 (ω)=ω 0 Gðt Þ sin ωt:dt, G00 (ω)=ω 0 Gðt Þ cos ωt:dt . Under linear strain-stress conditions, the shear stress response does reproduce the applied strain wave; i.e. it is a simple harmonic function of the applied wave. G0 > G00 indicates a solid-like behavior, while G00 > G0 indicates a viscous or flow behavior. Liquids and viscoelastic liquids are characterized by a flow behavior at low frequency, typically within 0.1–102 Hz. The surface interactions have a key role for the validity of the stress measurement. In a conventional measurement (millimeter thickness sample probed using aluminum, stainless steel, or glass substrates), the dynamic relaxation spectrum of polymers in the molten state exhibits, versus frequency, the typical Maxwell viscoelastic response (Ferry 1961). At low frequencies, the response is depicted by a ω-scaling decrease of the viscous modulus (G00 ) and a ω2-scaling of the elastic modulus (G0 ). G0 being negligible compared to the viscous modulus (Fig. 5a bottom), the flow behavior thus describes the frequency part where the viscous component becomes important. The interception of the two curves defines the terminal time τt, i.e., the largest time before the material enters in a flow regime. This characteristic time is interpreted as the longest molecular relaxation time (Rouse model).

256

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However, the apparent generic adequacy of the viscoelastic measurements to the Maxwell model should not hide the experimental divergences about the terminal time τt. Over large time scale relaxations (Boué et al. 1987; Wang 2006), unexplainable spectacular flow instabilities (spurt effect, “shark-skin” instability, shear-induced transitions) (Graham 1999; Pujolle-Robic and Noirez 2001), heterogeneous flows, and slippage at the boundaries (Mansard et al. 2014; Metivier et al. 2012) continue to feed a debate about the pertinence of the viscoelastic relaxation times to describe molecular dynamics over decades (Rault 1987; Litinov et al. 2013). The absence of chain deformation under steady-state shear flow in polymer melts at shear rates exceeding the inverse of the viscoelastic time (Watanabe et al. 2007; Noirez et al. 2009b – the photograph of Fig. 5 snapshotted at shear rates much larger than the inverse of the relaxation time shows optically a shear-thinning effect while there is no chain deformation) questions the entanglement/disentanglement concept. These observations point out the shortcomings of the macroscopic description in terms of molecular relaxation times while other works highlight nonlinear effects or highlight collective behavior and multiple intermolecular interactions. The recent consideration of the boundary conditions between surface and fluid interactions in rheology measurements has proven that the viscoelastic response is not universal but strongly influenced by the wetting or the anchoring conditions (Mendil 2006; Noirez, Baroni 2010). A totally different response, stronger and exhibiting a solid-like response, is obtained using total wetting conditions. The boundary conditions thus govern the efficiency of the transmission of the stress to the sample and play a major role in the quality of the measurements. It will be demonstrated that the optimization of the interaction between the surface and the material extends the dynamic relaxation spectrum and that the Maxwell viscoelastic response is actually only a part of a wider dynamic response.

Optimizing the Stress Transmission in Viscoelastic Measurements and Scanning the Submillimeter Scale Response Up-to-date progresses in rheology instrumentation allow the access to the measurement of the shear modulus with a high precision over six decades of magnitude. These improvements have considerably widened the frequency window and the access to very low stress moduli. Therefore, the detection of properties that would not been considered at the time of the first concepts of the viscoelasticity become accessible. These technical improvements can be conjugated to an optimization of the stress transmission by working on interfacial fluid/solid boundary conditions. The surface parameter is rarely taken into account. However, the validity of a dynamic relaxation experiment is entirely depending on the efficiency of the stress transmission which is ensured by the molecular contact forces between the sample and the surfaces. Because of symmetry reasons, interfacial energy differs from the 3D properties. The ideal boundary conditions correspond to lower the energy gap between two different media; i.e., the surface energy has to be as high as possible.

9

Probing Submillimeter Dynamics to Access Static Shear Elasticity from. . .

257

Fig. 5 (a) The upper figure schemes the viscosity η and the shear stress σ versus shear rate (flow curve). The rheofluidification zone (plateau) is interpreted by a shear-induced alignment of the chains. The bottom scheme shows the correspondence between the viscoelastic behavior and the flow curve with the interception of ω and ω2 scales of the viscoelastic curve defining the viscoelastic terminal time. (b) Conventional dynamic relaxation measurement of the macroscopic terminal behavior of the polybutadiene (PBD1,4). The reptation time is τrelax = 0.7  10
2 s at 26  C (cone-plate aluminum fixtures ARES rheometer). (c) Photograph of the PBD melt filling the Couette cylinder (gap thickness: 0.1 mm) and indicating that it is strongly shear stressed at 500 s
1 (shear-thinning regime). (d) Evolution of the radius of gyration of the PBD1,4 along the velocity (Rv) and the neutral axis (Rz) versus shear rate. The insets show the 2D neutron scattering patterns recorded at 30 s
1 and 770 s
1 (respectively, below and above the conventional terminal relaxation time) showing unchanged isotropic form factor of the polymer chain (Reprinted with permission from Noirez et al. 2009a)

A good criterion is the observation of a total wetting of the fluid onto the substrate (de Gennes et al. 2005 – Fig. 6b). The total wetting maximizes the molecular interactions to the surface. By filling voids, roughness, and asperities, it prevents the slippage. The discipline of rheology has grown ignoring the fluid/substrate boundary conditions on the dynamic measurement, treating in an equal way, wetting or non-wetting, hydrophilic or hydrophobic surfaces. Rheology supposes that the fluid wets adequately the metallic surfaces of the fixtures (generally made of aluminum or stainless steel). However, the determination of the contact angle shows that metallic substrates do not guaranty an optimal wetting (Fig. 6b, c).

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L. Noirez

a Medium surface energy: metal, glasses, PVC

b Very high surface energy: Functionalized surfaces, some metal oxides

Partial wetting

Total wetting µm

µm

c 70

Contact angle Q(°)

60 50 40 30 20 10 0

0

20

40

60 80 time (min)

Fig. 6 The wide majority of substrates (aluminum, stainless steel, glasses) exhibits partial wetting conditions. Only few fluid/substrate pairs fulfill the total wetting conditions. These conditions depend on the nature of both the fluid and the substrate. (a) The partial wetting is characterized by a finite Young contact angle and an incomplete wetting of the surface asperities (bottom profilometry scheme). (b) The zero-contact (macroscopic) angle of the total wetting ensures a complete wetting of the surface asperities (bottom profilometry scheme). (c) The contact angle is the angle where a liquid/air interface meets a solid surface. It quantifies the wettability of a solid surface  by a liquid and the slippage tendency using the Tolstoi’s relationship b  exp

σ 2 γ ð1
cos θÞ kT


1

where b is the slip length, σ a molecular constant, and γ the surface tension (Tolstoi 1952). The evolution of the contact angle versus time shows that the alumina substrate (red points: ) provides a total wetting (θ = 0), while the angle reaches a stationary value of 20 (partial wetting) for aluminum ( ), glass (□), and stainless steel ( ).

As a consequence of partial wetting conditions, the interfacial forces are reduced to molecules remaining in contact with the substrate, and the stress is not fully transmitted to the sample. Using the total wetting boundary conditions (alumina surfaces) and a conventional rheometer (Mendil et al. 2006; Baroni et al. 2005, 2010; Wang et al. 2007), a low-frequency solid-like response emerges at the submillimeter scale in a series of

G'

G''

10

1

2

10

w(rad/s)

0

0.1

PBuA Mw = 47 500 Tg = - 64°C

G''

G'

G(Pa)

10 -1 10

0

2

4

c 106

1

10

10

1

10

3

10

t(s) 5

2

10

w(rad/s)

PBuA Mw=40 000 T=26°C

10

0

0 e=0.025 mm

10 -1 10

10

10

0

10

2

10

2

1

-1

e=0.2 mm

10

0

10

e(mm)

10

5

G'

G''

4

6

10

10

Low frequency shear elasticity versus thickness

10

3

w (rad/s)

10

2

10

4

6

10

4

10 0.01

10

10

G'(Pa) 6 10

0

10 e=0.7 mm -1 1 10 10

10

2

10

4

6

10

G',G"(Pa)

Fig. 7 (a) Shear elasticity is progressively appearing at low frequency by decreasing the gap thickness (here a PBuA melt Mw = 40,000 Da studied at 90  C above the glass transition Tg =
64  C using wetting substrate). The probed thicknesses are e: 0.700 mm, 0.200 mm, 0.025 mm, respectively. (b) Low-frequency shear elasticity (plateau values at 10
1–10 rad/s) versus gap thickness. This solid-like (collective) response can be observed up to 1 mm in some liquids. (c) Relaxation modulus (G) versus time at constant shear strain (1%) (PBuA sample (Mw = 40,000 Da), wetting substrate, sample thickness 0.250 mm, room temperature). The modulus does not collapse with time in agreement with a solid-like behavior. (Reprinted with permission from Noirez et al. 2008)

b

G', G'' (Pa)

a G',G"(Pa)

9 Probing Submillimeter Dynamics to Access Static Shear Elasticity from. . . 259

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L. Noirez

polymer melts (Fig. 7). This observation was carried out on several ordinary polymers (polyacrylates, polymethacrylates, polybutadienes, etc.) tested from smaller to higher molecular weights (unentangled and entangled) up to high macroscopic scales (up to 0.5 mm gap thickness) and at about 100  C away from the glass transition and is in agreement with piezorheometer measurements carried at lower thicknesses (Derjaguin 1983; Derjaguin et al. 1989; Badmaev et al. 1983; Collin 2002).

From Surface-Induced Solidification to the Identification of the Submillimeter Shear Elasticity Low-frequency elastic behaviors have been reported at several occasions. Since 1991, Granick and coworkers (Hu et al. 1991; Demirel and Granick 2001; Zhu and Granick 2004) measured, at the molecular scale or a multiple of that, a solid-like component at nanoscales by SFA (Surface Force Apparatus) using mica surfaces. Close to a surface, these results were generally interpreted as resulting from surface-induced effects. The disappearance of the solid-like response at larger thicknesses is usually interpreted by a transition from surface to bulk properties. However the Mica (muscovite) surface does not provide strong liquid-substrate interactions (partial wetting) with organic fluids. Such boundary conditions are similar to those for conventional viscoelastic measurements. The very first analysis in terms of physical property is due to Derjaguin (Badmaev et al. 1983; Derjaguin et al. 1989) revealing in different fluids including polymers and liquid water, an elastic response at the micron scale that is interpreted as an intrinsic property generated by intermolecular interactions. At a larger scale, Collin et al. reported, using treated glass surfaces and small strains delivered by a piezorheometer, on a gel-like response up to 50 μm thicknesses in low molecular weight polystyrene melts that was interpreted as a reminiscence of the glass transition, i.e., clusters of finite size (Collin and Martinoty 2002). Since 2006 the use of optimized fluid/ substrate wetting conditions has facilitated the measurement of the low shear elasticity, generalizing its identification to ordinary liquids and confirming the probable origin due to intermolecular forces. The wetting protocol opens an easier access to the solid-like response usually hidden in conventional measurements of viscoelastic fluids. Figure 8 illustrates the response obtained using full wetting conditions for an H-bond oligomer poly(propylene glycol) (PPG4000) at room temperature. The flow behavior obtained conventionally on metallic substrate is here replaced by an elastic response (G0 and G00 nearly independent of the frequency and with G0 > > G00 ). The signal analysis indicates that the shear stress wave is superposed to the strain wave, confirming the instant transmission of the stress characterizing a solid-like behavior. This result is coherent with the conclusions in terms of elastic contribution carried out on the same liquid on the basis of the dynamic of the capillary waves (Chushkin et al. 2008).

Probing Submillimeter Dynamics to Access Static Shear Elasticity from. . .

a

3

10

2

10

1

10

b

G',G"(Pa)

0

CH3

O

HO C

CH3

O*

OH n

CH3

261

O CH3

G' G"

PPG 4000 T=+5°C, Tg=-75°C e=0.075mm

10

0

w(rad/s) 10

PPG 4000 e=0.075mm T= +5°C

Shear stress wave

10

Shear strain wave

9

1

Fig. 8 (a) Frequency dependence of the elastic G0 (ω): and viscous moduli G00 (ω): , measured for a glass former liquid (poly(propylene glycol) – PPG4000, MW = 4000 Da, Tg =
75  C) at T = +5  C (0.075 mm gap thickness – alumina plate-plate fixtures). (b) Superposition of the strain (green points) and the stress (red points) waves highlighting the instant response of the liquid

Generalization of the Submillimeter Shear Elasticity to Fluids and Liquids The low-frequency shear elastic regime is observed away from the glass transition (above a hundred degrees), on typically 0.025–0.500 mm sample thicknesses, thus in a liquid state where Ge is supposed to be zero. Using improved liquid/ substrate boundary conditions (i.e. high energy surfaces like the alumina to provide total wetting), the observation of a terminal elastic response has been also reported in glass formers ((PPG-4000; see Fig. 8), o-terphenyl (Noirez et al. 2011), glycerol (Noirez and Baroni 2010), and ordinary alkanes (Noirez et al. 2012)) and even on liquid water. The low frequency shear elasticity cannot be interpreted by entanglement effects and the vicinity of the glass or of a crystalline transition but is due to intermolecular interactions ensuring the liquid cohesion. Being measured at several tenths or hundredths of millimeter and reaching several thousands Pascals (polymer melts), the low-frequency shear elasticity cannot be interpreted by a surface-induced solidification. A solid-like response is even reported up to the millimeter length scale in the isotropic phase of a liquid-crystalline polymer (Noirez 2005), while the molecular dimensions are less than 100 Å. The strong anchoring of the liquid crystal molecules reinforces the boundary contacts between the fluid and the substrate whereby the stress is transmitted and facilitates the measurement (Fig. 9). Gallani et al. observed, using the molecular displacement strain produced by the piezorheometer, an “abnormal” viscoelastic behavior on tens micron sample thicknesses in the isotropic phase of a liquid-crystalline polymer substrates using a surface treatment (Gallani et al. 1994; Martinoty et al. 1999).

262

L. Noirez Shear vs Normal forces

b 0,1

0,05

0

0

–0,1 0

–0,05 10

20

30

40

103

G′,G′′(Pa)

Normal force (volts)

Shear Strain Force (volts)

a

t(s)

c 103

G′,G′′(Pa)

103 G′,G′′(Pa)

Noisy signals

G′

G′ 2

10

101

102

G ′′ G′

101

100

102 G ′′

101

100 10–1

100 ω(rad/s)

10–1

100 ω(rad/s)

100

G ′′

Shear elastic plateau

10–1

100 ω(rad/s) time

Fig. 9 (a) Photograph of the alumina fixtures (surface roughness 1, neutral λ = 1, and non-aligning λ < 1 LCPs are used (de Andrade Lima and Rey 2004a-d; Martins 2001) (see Table 2): (i) Aligning LCPs: (1) PSi4 (poly[(2,3,5,6tetradeuterio-4methoxyphentl-40 -butanoxybenzoate)-methylsiloxane]), (2) AZA9 (poly(4,40 dioxy-2, 20 -dimethylazoxybenzene-dodeccanediyl)), and (3) DDA9 (poly(4,40 -dioxy-2, 20 dimethylazoxybenzene-dodeccanediyl)). (ii) Neutral LCP: (4) TPB10 (poly[1,10decylene-1-(4hydroxy-40 -biphenylyl)-2-(4-hydroxyphenyl) butane]). (iii) Non-aligning LCPs: (5) PBLG (poly(γbenzyl-L-glutamate)) 17% in m-cresol, (6) PPTA 8.8% (poly ( p-phenylene terephthalamide)) in SO4D2, and (7) PBLG 12% in m-cresol. Imposing pressure oscillations on the NLCs will produce spatially nonhomogeneous director oscillations. Since NLCs are viscoelastic materials, the director oscillations will not be in-phase with the applied pressure drop. Thus, the total director angle θ(r, t, ω) is given by the sum of the following in-phase and out-phase components: θðr,tÞ ¼ θi ðr,tÞ sin ðωtÞ þ θ0 ðr,tÞ cos ðωtÞ

(30)

Since the director field n is coupled to the velocity field v, imposing an oscillatory pressure drop to the NLC will produce a velocity field with in-phase and out-ofphase components (de Andrade Lima and Rey 2004a-e). Thus, the total dimensionless velocity field v(r, t, ω) is given by the sum of the following in-phase and out-phase c components: vðr,t,ωÞ ¼ vi ðr,ωÞ sin ðωtÞ þ v0 ðr,ωÞ cos ðωtÞ

(31)

Using the in-phase and out-phase dimensionless components in Eqs. 30 and 31, the following expressions were obtained for the dimensionless storage modulus G0 , loss modulus G00 , and loss tangent tan δ = G00 /G0 (de Andrade Lima and Rey 2004a-e):

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Liquid Crystalline Polymers: Structure and Dynamics

299

Table 2 Parameter values (Martins 2001) 1

2

3

4

Set T(K) α1 α2 α3 α4 α5

PSi4 348 960 2000 127.0 3066 1534

AZAA9 393 1320 1595 25.00 208.9 1470

DDA9 394 162.0 170.0 2.000 16.01 162.0

TPB10 360 1.793  105 1.810  105 0.000 7.838  105 1. 752  105

5 PBLG 17% in m-cresol 302 960 2000 127.0 3066 1534

α6 λ (Reactive parameter)

593.2 1.1356

149.9 1.0318

10.01 1.0238

5.838  103 1.0000

593.2 0.9940

6 PPTA 8.8% In SO4D2 300 1.177  106 2.136  106 6.972  104 2.415  105 1. 767  106 2.992  105 0.9368

M1 ½ωF 2  1 

F1 1 2 1þ þ FM2  ωM M   ωF 2 þω M2 M  G00 ðωÞ ¼ 

F1 1 2 1þ þ FM2  ωM M G0 ðωÞ ¼ 

tan ðδÞ ¼

M þ F 1 ð ωÞ 1 F2 ðωÞ  ω

7 PBLG 12% in m-cresol 302 193.0 390.0 29.00 59.83 291.6 69.43 0.8616

(32)

(33)

(34)

where de functions F1(ω) and F2(ω) pffiffiffiffi pffiffiffiffi pffiffiffiffi pffiffiffiffi pffiffiffiffi pffiffiffiffi ber1 ωðbei0 ω  ber0 ωÞ  bei1 ωðber0 ω þ bei0 ωÞ pffiffiffiffiffiffi F1 ðωÞ ¼ pffiffiffiffi pffiffiffiffi

2 2ω ber21 ω þ bei21 ω pffiffiffiffi pffiffiffiffi pffiffiffiffi pffiffiffiffi pffiffiffiffi pffiffiffiffi bei1 ωðber0 ω  bei0 ωÞ  ber1 ωðbei0 ω þ ber0 ωÞ pffiffiffiffiffiffi 2 pffiffiffiffi

F2 ðωÞ ¼ p ffiffiffi ffi 2 2ω ber1 ω þ bei21 ω   3π π v þ k cos 1 x2kþv X 4 2 F2 ðωÞ ¼ berv ðxÞ ¼ k!Γðk þ 1 þ vÞ 2 k¼0   3π π v þ k  2kþv 1 sin X x 4 2 F2 ðωÞ ¼ beiv ðxÞ ¼ k!Γðk þ 1 þ vÞ 2 k¼0

(35)

(36)

(37)

(38)

300

A. D. Rey et al.

The functions beiν(x) and berν(x) are the Kelvin functions of order ν (de Andrade Lima and Rey 2004a-e). Note that these results have the same pattern of the results presented for simple shear flow of NLCs. The main viscoelastic parameter ratios that the dimensionless storage modulus G0 , the loss modulus G00 , and the loss tangent tanδ are respectively M1, M2 and M given by:  2 α3 M1 ¼ 8 ¼ 2 ð 1  λÞ 2 γ1 η1 γ1

(40)

1 η1 γ1 8 α3 α3

(41)

M2 ¼ M¼

(39)

The storage modulus G0 increases with M1 since this ratio increases the viscous torques creating more elastic storage and less rotational dissipation; when λ = 1, viscous torques are absent and no elastic storage is possible G0 = 0. The loss modulus G00 increases with M2 since this ratio increases translation dissipation and is larger than rotational dissipation; when γ1 ! 1 rotational dissipation dominates G00 = 0. The ratio M is the product of the two dissipation torques ratios and can be rewritten in terms of λ as: M1 ¼

1 η1 1 2 γ 1 ð 1  λÞ 2

(42)

By using the known asymptotic behavior (de Andrade Lima and Rey 2004a-e) of the Kelvin functions, the frequency dependence of the viscoelastic moduli for α3 6¼ 0 is as follows: the loss modulus is always greater than the storage modulus, the low frequency (terminal) regime is classic of a viscous fluid, and characteristic slopes are:



a ω ! 0,G0  ω2 ,G00  ω; b ω ! 1,G0  ω1=2 ,G00  ω

(43)

It follows from Eq. 43, the behavior corresponds to a Newtonian material and G0 = 0. In addition, at frequencies larger than resonance, the dependency of material properties simplifies; resonance is found with ωr = 18.6522, and the large frequency regime results in this section hold for ω > 10 ωr. The factorization of material properties from frequency dependent functions is obtained by direct order of magnitude analysis. At frequencies larger than the resonance ωr, the terms F1/M and F2/ M – 1/ωM in Eqs. 32 and 33 are very small compared with one; in addition, in the numerator of Eq. 33 F2/M is less than the unity; therefore, the asymptotic expressions for the viscoelastic functions for frequencies larger than ωr are: G0 ¼ ωF2  1 ¼ Φ0 ðωÞ M1

(44)

10

Liquid Crystalline Polymers: Structure and Dynamics

G00 ¼ ω ¼ Φ00 ðωÞ M2   tanδ 1 1 ¼ F 2 ð ωÞ  ¼ Φ00 ðωÞ M ω

301

(45)

(46)

These equations show the origin of the vertical (amplitude) scaling of all data sets. The horizontal scaling (frequency) is obtained by plotting the dimensionless storage modulus G0 , and loss modulus G00 , and loss tangent (tanδ) as functions of “ω.” Figure 12 shows the scaled storage modulus (G0 /M1) and scaled loss modulus (G00 /M2) as a function of the dimensionless frequency ω for the data sets. The figure shows for storage modulus that a collapse of the curves is essentially perfect for ω > 30 ωr. For TPB10, the storage modulus is zero. The figure also shows a collapse of the curves of loss modulus of essentially perfect for ω > 16 ωr. Figure 13 shows the scaled loss tangent (tanδ/M) as a function of the dimensionless frequency ω for the data sets. These results show a collapse of the plots is almost perfect especially to high frequencies. The rheological responses of defect-free liquid crystal polymers subjected to small amplitude oscillatory pressure-driven Poiseuille flow show superposition and universality. In general, the degree of superposition is almost perfect at frequencies above resonance. Although liquid crystals are usually differentiated into shear aligning and non-aligning materials, the results presented here shows that under certain flow conditions, linear viscoelasticity only depends on the dimensionless factor (1  λ)2, and hence rheological equivalence between aligning and non-aligning materials exists if λNA = 2  λA. Moreover, the resonance is also shown to be independent of flow alignment.

Flow Birefringence of Biological Liquid Crystals Flow-birefringence in the isotropic phase is observed in all types of liquid crystals including LCPs and is a manifestation of flow-induced orientation and ordering due to viscous torques acting on the anisodiametric mesogens. The flow birefringence due to transient shear flow was detected using optical transmittance which is proportional to the square of the order parameter S produced by the shear flow (Rey 2009, 2010; Rey et al. 2014). By fitting experimental data with the LdG results, the rotational diffusivity can be estimated. Here, the LdG model for flowbirefringence under steady flow is used to illustrate the estimation of rheological parameters from the experimental signal (Rey 2009, 2010; Rey et al. 2014). Using a planar director field, simple shear, and steady uniaxial nematodynamics (Eq. 13), the following relation between flow-alignment θal and S is predicted (Rey 2009, 2010; Rey et al. 2014):

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A. D. Rey et al.

~

SCALED DIMENSIONLESS LOSS MODULUS (G"/M2)

Fig. 12 Scaled dimensionless storage (G0 /M1) and loss (G00 /M2) moduli as a function of the dimensionless frequency ω for Psi4, AZA9, DDA9, PBLG 17%, PPTA 8.8%, and PBLG 12%

1000000

100000

10000

1) PSi4 2) AZA9 3) DDA9 5) PBLG 17% in m-cresol 6) PPTA 8.8% in SO4D2 7) PBLG 12% in m-cresol 2

1000

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1,2,3,5 7,6

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0.1 0.0001

0.001

0.01

0.1

1

10

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SCALED DIMENSIONLESS STORAGE MODULUS (G'/M1)

Fig. 13 Scaled loss tangent tanδ/M as a function of the dimensionless frequency ω for Psi4, AZA9, DDA9, PBLG 17%, PPTA 8.8%, and PBLG 12%

1 6S

, 1 þ cos 2θal ¼ ¼ λ β 4 þ 2S  S2



ð3  UÞS  US2 þ 2US3 De

2

!  1 λ2 ¼ 0

(47) where for simplicity, the equality 18η = ς is used to relate translational (η) and rotational (ς) viscosities (see for example Eq. 30 of Rey 2010). Figure 14a shows the computed scalar order parameter S and the alignment angle as a function of De for steady shear of an isotropic solution, for β = 0.95 and several values of U; the shapes

Order Parameter, S

a

303

1

50

0.8

40 θ@U=2.6 θ@U=2 θ@U=1 S@U=1 S@U=2 S@U=2.6

0.6

0.4

0.2

0

b

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10

0

0 15

5 10 Deborah number, De

1 S @ β=0.5 S @ β=0.7 S @ β=0.95

0.8 Order Parameter, S

30

Alignment angle, θ

Liquid Crystalline Polymers: Structure and Dynamics

50

θ@ β=0.5 θ@ β=0.7 θ@ β=0.95

40

0.6

30

0.4

20

0.2

10

0 0

1

2

3

4 5 6 7 Deborah number, De

8

9

Alignment angle, θ

10

0 15

Fig. 14 Flow-induced birefringence (S > 0, 0 < θal < π/4) as a function of De of a sheared isotropic solution of rods, (a) different concentration of rods (U) and β = 0.95; increasing U increases the initial slope dS/dDe. (b) Same but for U = 2 and different β values increasing βincreases the plateau value of S. (Adapted from Rey 2010)

of S curves are consistent with experiments (Rey 2009, 2010; Rey et al. 2014). Flowinduced birefringence is due to increasing S and decreasing θal as shear rate increases. From Fig. 14a and through Eq. 47, the limits for low and high Deborah number can be obtained: De 1,So
0,θal
π=4; De  1,S1 qffiffiffiffiffiffiffiffiffiffiffiffiffiffiffiffiffiffiffiffiffiffiffiffiffiffiffiffiffiffi
ð1  3=βÞ þ ð1  3=βÞ2 þ 4,θal
0

(48)

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hence, S(θal) increases (or decreases) monotonically from zero (π/4) to saturate at a β-dependent plateau S1(0). Figure 14b shows that flow-birefringence increases with increasing β (or rheological shape anisotropy); this prediction can be used to estimate β and then λ, the reactive parameter through Eq. 47.

Patterns and Textures Banded Patterns Banded patterns normal to the flow direction after cessation of flow form for shear rates up to 50 s1; the pattern formation time is proportional to the square of the applied shear, and the pattern coarsening rate at sufficiently low pre-shear is well described by a diffusive process. The pattern is associated with relaxation of elastic energy. For shear rates greater than 50 s1, the material flow-aligns. These features were found to be similar to those found in non-aligning lyotropic nematic polymers (Roux et al. 1995). Using given flow kinematics and a 1D spatial description, a full rheological phase diagram is obtained based on the LdG for non-aligning nematics in terms of the length scale ratio R and the Ericksen number Er. The left schematic on Fig. 15 shows the rheological phase diagram in terms of the eight director modes across the cell thickness. For low Er, elasticity prevails, and tumbling is arrested. The dotted parabola corresponding to small Er, containing the out-of-plane modes, regions 3–5 display multistability. The main features of these flow modes are summarized in the following paragraph (Fig. 15 left): 1. In-Plane elastic-driven steady state (EE): The steady state of this planar mode arises due to the long rate order elasticity stored in the deformed tensor order parameter field. In this planar mode, there is no orientation boundary layer behavior because there is no flow-alignment in the bulk region 2. In-plane tumbling wagging composite state (IT): In this time-dependent planar mode, the director dynamic in the bulk region is rotational and in the boundary layer it is oscillatory. The boundary between the bulk region and each boundary layer is characterized by the periodic appearance of the abnormal nematic state, which is characterized by two equal eigenvalues of the tensor order parameter (i.e., μk = μr > μ1) and follows a smoothly defect-free transition from the rotation bulk region to the fixed director anchoring at the surfaces by a director resetting mechanism. The insert in Fig. 15 is discussed in full detail below and shows schematics of the existing stable solutions of Eq. 2 3. In-plane wagging state (IW): In this plane mode, the director dynamics over the entire flow geometry is periodic oscillatory with an amplitude decreasing from a maximum at the centerline to zero at the two bounding surfaces

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Fig. 15 (Left) Rheological phase diagram as a function of the ratio of short–long range elasticity (R) and the ratio of viscous flow to long range elasticity effects (Er), and corresponding director configurations. There are eight flow regimes in the parametric space Er and R and nine flow modes. Lines represent flow regime transitions. The dotted line shows the transition between in-plane and out-of-plane modes. The arrows represent the director, and empty circles are the abnormal nematic state. (Right) Shows flow modes observation probability as a function of the Ericksen number. The figure presents the concepts of flow regime transitions in terms of the change in the observation probability with increasing Er. (Adapted from Tsuji and Rey 1998)

4. In-plane viscous-driven steady state (IV): In this plane mode, the director profile shows a flow-aligning bulk region and two boundary layers. On traversing the boundary, the director rotates form the aligning angle to the flow direction at the walls 5. Out-of-plane elastic-driven steady state with achiral structure (OEA): In this nonplanar mode, the director shows steady twist structures, and the twist angle profiles are symmetric with respect to the centerline. The steady state arises due to the long-range order elasticity. Similar solutions are presented by the LeslieEricksen solutions. Following from the bottom to top bounding surface, the net director twist rotation is null 6. Out-of-plane elastic-driven steady state with chiral structure (OEC[n]): In this nonplanar mode, the director shows steady twist structure, with nπ (n = 1, 2) radian difference between the anchoring angles at the lower and upper bounding surfaces, but without the presence of defects or disclinations. The different anchoring conditions are smoothly connected by the chiral director structure. A similar OEC solutions is predicted by the LE equations

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7. Out-of-plane tumbling-wagging composite state with periodic chirality (OTP): In this nonplanar mode, the bulk director dynamics is planar and rotational, and in the two boundary layers, it is nonplanar and rotational, and in the two boundary layers, it is nonplanar oscillatory. The spatial profiles of the periodic director motion are anti-symmetric. The director field exhibits periodic chirality, that is, after a cycle of 2π rotation of the bulk director, the system periodically recovers the spatially homogeneous director configurations (i.e., n ffi (1, 0, 0)) for 0  y  1 8. Out-of plane tumbling-wagging composite state withπ chiral structure (OTC). The director dynamics is in-plane rotational in the bulk region and out-of plane oscillatory in the boundary layers, the directors at the upper and lower bounding surfaces have opposite directions, and the system never recovers to the spatial homogeneous director configuration. Figure 15 is a schematic of the rheological phase diagram given in terms of R and Er, clearly indicating the parametric regions where the four planar modes and the five nonplanar modes are predicted. Finally, observation probabilities of all flow modes for all the flow regimes are shown in Fig. 15b where the probabilities are plotted for each flow mode, and thus for any Ericksen number, the sum of the probabilities is 1. For example, for region I, the probability of IE mode is 1 and others are zero, and for region 4 the probability of OEA and OTP flow modes are almost 0.5, and the others are zero. Summarizing this subsection, the results presented here predict extensive multi-stabile phenomena, involving planar, chiral achiral, steady, and time periodic nodes. The range and richness of the multistability is due to the presence of the two compatibilization mechanisms predicted by the complete theory (Rey and Tsuji 1998; Rey and HerreraValencia 2012).

Banded Textures After Cessation of Shear Banded texture predictions during flow and after cessation of flow have been compared with experimental data using LE and LdG models. The consistency between the two models’ predictions have been established (Rey and Tsuji 1998). In the LE model of banded textures during flow, the pattern formation is driven by the OP mode that nucleates a periodic array of elliptical splay-twist-bend inversion wall in the velocity/velocity gradient plane with a wave-length close to the shear cell thickness. In the LdG model of banded textures, the nucleation and growth of OP modes give rise to a heterogeneous nonplanar orientation field that relaxes through the formation of a periodic texture. Figure 16 shows the out-of-plane component profile after cessation of flow, for R = 0 at the following dimensionless times: (a) t = 2, (b) 4, (c) 6, (d) 8, (e) 12, (f) 16, and (g) 20. The small source of the out-of-plane component near the surface is connected across the bulk region, and the banded texture is formed with almost the same director configuration as that during flow. Then, the texture relaxes through the shrinking of the director out-of-plane region.

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1 (a) t*=2

(b) t*=4 0.5 (c) t*=6

(d) t*=8

0

(e) t*=12

(f) t*=16

(g) t*=20

Fig. 16 Director out-of- plane component nz profile after cessation of flow, for R = 0 at the following dimensionless times: (a) t = 2, (b) 4, (c) 6, (d) 8, (e) 12, (f) 16, and (g) 20. Er = 100 of the shear flow is applied until t = 100, and thus the result at t = 100 for Er = 1000 corresponds to that at t = 0 for Er = 0 (no flow). The color scale on the right indicates the correspondence of the magnitude of nz and the shown colors. The relaxation of stored elastic energy drives the formation of a spatial pattern. (Adapted from Tsuji and Rey 1998)

Biological Liquid Crystalline Polymer Processes The objective of this section is to show that once rheological characterization of biological liquid crystals is carried out, such knowledge can be used to explore processes reported in vitro as is the case of film casting of collagen type I solutions and in vivo as is the case of silk spinning.

Film Casting of Cholesteric Collagen Solutions Dilute collagen solutions have been processed into thin films as shown in Fig. 17 to form defect-free cholesteric films with other transport processes as the evaporation of solvent (Rey 2010). A key feature of the chiral film is the nano-scale surface

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Fluid velocity, VF

l

Fig. 17 Film casting of collagen solutions for tissue engineering (ex.: artificial cornea). The surface of the chiral film has nano-wrinkling features due to the chirality of collagen

wrinkling whose wave length is equal to the pitch of the collagen helix. This wrinkling has been shown to be beneficial in tissue engineering applications and its origin emerges when considering the anisotropic component of the interfacial tension. Since capillary tension depended on the relative angle between the collagen rods and the surface normal, this angle is a periodic function due to the chirality of collagen and hence nano-wrinkling of wave length λ arises.

Silk Spinning The silk spinning process includes synthesis of water soluble (random coil) proteins, storage and liquid crystal self-assembly, chemo-rheological processing, and solidification into a water-insoluble nanocomposite fiber containing β-sheet crystallites (Rey 2010 and references therein). The intermediate stages of self-assembly and precursor processing are responsible for the orientational structuring and alignment (Rey and Herrera-Valencia 2012; Rey 2010). The chemo-rheological processing includes flow in a coated double exponential capillary, dehydration, acidification, and metallic cation gradients; the last two effects are believed to increase the conformational change response to shear and extensional flow (Rey and HerreraValencia 2012). Experimental evidence shows that the intermediate phase between the protein solution and the insoluble silk fiber is a liquid crystal that appears in the sac and first section of the duct components of the spinning process. Figure 18 summarizes the observations of textural transitions and geometry changes along the duct of a silkworm; the similarities and correspondence with the major ampullate silk gland duct of the orb web spider Nephila is discussed by Rey and Herrera-Valencia (2012), Rey (2010), and references therein. The central region in the figure shows a texture sequence along the duct, where the rods denote the average orientation where the geometry was extracted and adapted from Rey and Herrera-Valencia (2012), Rey (2010), and references therein. In the entry section, the nematic phase adopts an escape (E) texture with the rods along the axis at the centerline and normal to the sericin coated walls. When the duct diameter reaches close to 80 μm, there is a texture transformation from escape to a point defect texture (PD) known as cellular optical texture. This typical nematic texture under capillary confinement is a 1D defect lattice, consisting of alternating 1 point defects, separated by a distance

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Fig. 18 Texture transformation in the duct of a silkworm. The average orientation changes along the duct section of the silkworm spinning gland coated with sericin. The three textures are: escape (E), point defect (PD), and isotropic (I) (Rey and Herrera-Valencia 2012; Rey 2010; and references therein)

close to the duct diameter. When the duct diameter reaches 60 μm, the cellular optical texture is replaced by an isotropic solution that lacks birefringence despite the presence of capillary flow. The predictions of nematodynamics show consistency between LCP models and silk spinning observations shown in Fig. 18 (Rey and Herrera-Valencia 2012; Rey 2010; and references therein) and reveal a structuring mechanism in silk spinning.

Conclusions A comprehensive review of the theory and computer simulations of lyotropic and thermotropic liquid crystalline polymers’ flow processes was presented based on the classical Leslie-Ericksen and Landau – de Gennes models. Emphasis was made on the structure and dynamics of these materials, including defects, texturing, and shear rheology. Strong similarities and analogous structural properties and dynamics behavior between lyotropic and thermotropic including flow-birefringence, banded textures under magnetic fields, thermo-rheological responses across IN transitions, and shear-rheology scaling were discussed in detail. Although not included in this chapter, other emerging applications analyzed with the presented theory are discussed elsewhere (Rey et al. 2014; Rey and Herrera-Valencia 2012; Rey 2007, 2008a, b, 2009; 2010). The chapter provides an integrated framework to characterize flow-processes of liquid crystal polymers using accurate models. Acknowledgments ADR was supported by the Natural Science and Engineering Research Council of Canada (NSERC), Compute Canada, Calcul Quebec, and McGill University. EEHV gratefully acknowledges financial support of PASPA for the sabbatical research at the Chemical Engineering Department, McGill University and PAPIIT, PAPIME projects IN115919, PE116519 from DGAPA/UNAM, respectively. OFAG gratefully acknowledges financial support from CONACYT-MEXICO (Doctoral Grant no 313480). This research is dedicated to the memory of my beloved father Emilio Herrera Caballero. Thank you so much for your help Dr. Edtson Emilio Herrera Valencia

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Ericksen JL (1969) A boundary-layer effect in viscometry of liquid crystals. Trans Soc Rheol 13:9–15 Farhoudi Y, Rey AD (1993a) Shear flow of nematics polymers. I. Orienting modes, bifurcations, and steady state rheological predictions. J Rheol 37:289–314 Farhoudi Y, Rey AD (1993b) Shear flow of nematic polymers. II Stationary regimes and start-up dynamics. J Non-Newtonian Fluid Mech 49:175–204 Farhoudi Y, Rey AD (1993c) Ordering effects in shear flows of discotic polymers. Rheol Acta 32:207–217 Fratzl P (2003) Cellulose and collagen: from fibres to tissues. Curr Opin Colloid Interface Sci 8:32 Fratzl P, Giraud-Guille MM (2011) Hierarchy in natural materials. In: Su BL, Sanchez C, Yang XY (eds) Hierarchically Structured Porous Materials. Wiley-VCH, pp 29–39 Giraud-Guille MM (1998) Plywood structures in natures. Curr Opin Solid State Mater Sci 3:221–227 Giraud-Guille MM (2005) Bone matrix-like assemblies of collagen: from liquid crystals to gels to biomimetic materials. Micron 36:602–608 Grecov D, Rey AD (2003a) Shear-induced textural transitions in flow aligning liquid crystals polymers. Phys Rev E 68:061704–061724 Grecov D, Rey AD (2003b) Theoretical computational rheology for discotic nematic liquid crystals. Mol Cryst Liq Cryst 39:157–194 Grecov D, Rey AD (2003c) Transient shear rheology of discotic mesophases. Rheol Acta 42:590–604 Grecov D, Rey AD (2004) Impact of textures on stress growth of thermotropic liquid crystals subjected to step-shear. Rheol Acta 44:135–149 Grecov D, Rey AD (2006) Texture control strategies for flow-aligning, liquid crystal polymers. J Non-Newtonian Fluid Mech 139:197–208 Gupta G, Rey AD (2005a) Texture rules for concentrated filled nematics. Phys Rev Lett 95:127805/ 1-4 Gupta G, Rey AD (2005b) Texture modeling in carbon-carbon composites based on mesophase precursor matrices. Carbon 437:1400–1406 Han WH, Rey AD (1994a) Orientation symmetry breakings in shear liquid-crystals. Phys Rev E 50:1688–1691 Han WH, Rey AD (1994b) Dynamic simulations of shear-flow-induced chirality and twisted textures in a nematic polymer. Phys Rev E 49:597–613 Han WH, Rey AD (1995a) Theory and simulation of optical banded textures of nematic polymers during shear flow. Macromolecules 28:8401–8405 Han WH, Rey AD (1995b) Simulation and validation of temperature effects on the nematorheology of aligning and non-aligning liquid crystals. J Rheol 39:301–323 Herrera-Valencia EE, Rey AD (2014) Actuation of flexoelectric membranes in viscoelastic fluids with applications to outer hair cells. Phil Trans R Soc A 372:2013.0369 Ikoma T, Kobayashi H, Tanaka J, Walsh D, Mann S (2003) Microstructure, mechanical and biomimetic properties of fish scales from Pagrus major. J Struct Biol 142:327–333 Kirkwood JE, Fuller GG (2009) Liquid crystalline collagen: a self-assembled morphology for the orientation of mammalian cells. Langmuir 25:3200–3296 Knight DP, Feng D (1994) Interaction of collagen with hydrophobic protein granules in the egg capsule of the dog fish Scyliorhinus canicula. Tissue Cell 26:155–167 Kundu S, Grecov D, Ogale A, Rey AD (2009) Shear flow induced microstructure of a synthetic mesophase pitch. J Rheol 53:85–113 Kupchinov B, Ermakov S, Rodnenkov V, Bobrysheva S, Beloenko E, Kestelman V (1993) Role of liquid crystals in the lubrication of living joints. Smart Mater Struct 2:7–12 Larson R (1999) The structure and rheology of complex fluids. Oxford University Press, New York Larson RG, Doi M (1991) Mesoscopic domain theory for textured liquid crystalline polymers. J Rheol 35:539–563

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Li J, Revol JF, Marchessault RH (1996) Rheological properties of aqueous suspensions of chitin crystallites. J Colloid Interface Sci 183:365–373 Livolant F, Bouligand Y (1978) Double helical arrangement of spread dinoflagellate chromosomes. Chromosoma 68:21–44 Livolant F (1991) Ordered phases of DNA in vivo and in vitro. Physica A Stat Mech Appl 176:117–137 Lydon J (2006) Microtubules: nature smartest mesogens-a liquid crystal model for cell division. Liq Cryst Today 15:1–10 Martins AF (2001) Measurement of viscoelastic coefficients for nematic mesophases using magnetic resonance. In: Dunmur DA, Fukuda A, Luckhurst G (eds) Physical properties of liquid crystals: nematics. IEE Publishing, London, pp 405–413 Miller AF, Donald AM (2003) Imaging of the isotropic/anisotropic surfaces of aqueous cellulose suspensions using environmental scanning electron microscopy. Biomacromolecules 4:510–517 Murugesan YK, Rey AD (2010) Structure and rheology of fiber–laden membranes via integration of nematodynamics and membranodynamics. J Non-Newtonian Fluid Mech 165:32–44 Murugesan YK, Rey AD (2011) Mechanics of fiber-laden membranes. Contin Mech Thermodyn 23:45–61 Murugesan YK, Pasini D, Rey AD (2011) Microfibril organization modes in plant cell walls of variable curvature: a model system for two dimensional anisotropic soft matter. Soft Matter 7:7078–7093 Neville AC (1993) Biology of fibrous composites: development beyond the cell membrane. Cambridge University Press, Cambridge Neville AC, Luke BM (1971) A biological system producing a self-assembling cholesteric protein liquid crystal. J Cell Sci 8:93–109 Petrov AG (2002) Flexoelectricity of model and living membranes. Biochim Biophys Acta Biomembr 1561:1–25 Revol JF, Bradford H, Giasson J, Marchessault R, Gray D (1992) Helicoidal self-ordering of cellulose microfibrils in aqueous suspension. Int J Biol Macromol 14:170–172 Rey AD (1993a) Analysis of shear-flow effects on liquid-crystalline textures. Mol Cryst Liq Cryst 225:313–335 Rey AD (1993b) Rheological prediction of a transversely isotropic fluid model with extensible microstructure. Rheol Acta 32:447–456 Rey AD, Tsuji T (1988) Recent advances in theoretical liquid crystals in rheology. Macromol Theory Simul 7:623–639 Rey AD, Denn MM (2002) Dynamical phenomena in liquid-crystalline materials. Annu Rev Fluid Mech 34:233–266 Rey AD (2005) Mechanics of soft solid-liquid crystals interfaces. Phys Rev E 72:011706 Rey AD (2006) Liquid crystals model of membrane flexoelectricity. Virtual J Biol Phys Res 12:011710 Rey AD (2007) Capillary models for liquid crystals fibers, membranes, films and drops. Soft Matter 2:1349–1368 Rey AD (2008a) Nonlinear actuator model for flexoelectric membranes. Int J Des Nat Ecodyn 3:28–38 Rey AD (2008b) Linear viscoelastic model for bending and torsional modes in fluid membranes. Rheol Acta 47:861–871 Rey AD (2009) Flow and texture and modeling of liquid crystalline materials. Rheol Rev 2008:71–135 Rey AD (2010) Liquid crystals model of biological materials and processes. Soft Matter 6:3402–3429 Rey AD, Herrera-Valencia EE (2012) Rheological theory and simulation of surfactant nematic liquid crystals. In: Garti N, Somasundaran P, Mezzenga R (eds) Self-assembled supramolecular architectures: lyotropic liquid crystals. John Wiley & Sons Inc. Hoboken, New Jersey USA

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Rey AD, Herrera-Valencia EE, Murugesan YK (2014) Structure and dynamics of biological liquid crystals. Liq Cryst 41:430–451 Rey AD, Tsuji T (1998) Recent advances in theoretical liquid crystals rheology. Macromol Theory Simul 7:623–639 Roland J, Reis D, Vian B, Satiat-Jeunemaitre B, Mosiniak M (1987) Morphogenesis of plant cell walls at the supramolecular level: internal geometry and versatility of helicoidal expression. Protoplasma 140:75–91 Roux DC, Berret JF, Porte G, Peuvrel-Disdier E, Lindner P (1995) Shear-induced orientation and textures of nematic living polymers. Macromolecules 28:1681–1687 Sharma V, Crne M, Park JO, Srinivasarao M (2009) Structural origin of circularly polarized iridescence in jeweled beetles. Science 325:449–451 Soulé ER, Abukhdeir NM, Rey AD (2009) Thermodynamics, transition dynamics, and texturing in polymer-dispersed liquid crystals with mesogens exhibiting a direct isotropic/smectic A transition. Macromolecules 42:9486–9497 Tsuji T, Rey AD (1997) Effect of long range order on sheared liquid crystalline materials. Part I: compatibility between tumbling behaviour and fixed anchoring. J Non-Newtonian Fluid Mech 73:127–152 Tsuji T, Rey AD (1998) Orientation mode selection mechanisms for sheared nematic liquid crystalline materials. Phys Rev E 57:5609–5625 Tsuji T, Rey AD (2000) Effect of long range order on sheared liquid crystalline materials, transition and rheological phase diagrams. Phys Rev E 62:8141–8151 Vollrath F, Knight DP (2001) Liquid crystalline spinning of spider silk. Nature 410:541–548 Willcox PJ, Gido SP, Muller W, Kaplan DL (1996) Evidence of a cholesteric liquid crystalline phase in natural silk spinning processes. Macromolecules 29:5106–5110 Wright DC, Mermin ND (1989) Crystalline liquids: the blue phases. Rev Mod Phys 61:385–432 Yan J, Rey AD (2003) Modeling elastic and viscous effects on the texture of ribbon-shaped carbonaceous mesophase fibers. Carbon 41:105–121

Part II LCP Functions and Applications

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Matti Knaapila, Roman Stepanyan, and Andrew P. Monkman

Contents Definition . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . Introduction . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . Conjugated Polymers as Liquid Crystalline Hairy-Rod Polymers . . . . . . . . . . . . . . . . . . . . . . . . . . . . . Concept . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . Thermotropic Behavior . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . Lyotropic Behavior . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . Alignment . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . Conjugated Polymers as Liquid Crystalline Hairy-Rod Supramolecules . . . . . . . . . . . . . . . . . . . . . . Thermotropic Behavior . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . Lyotropic Behavior . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . Challenge . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . Cross-References . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . References . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . .

318 318 320 320 321 323 326 330 330 332 333 336 337

M. Knaapila (*) Department of Physics, Technical University of Denmark, Kgs. Lyngby, Denmark e-mail: [email protected] R. Stepanyan Materials Science Centre, DSM Research, Geleen, The Netherlands e-mail: [email protected] A. P. Monkman Department of Physics, University of Durham, Durham, UK e-mail: [email protected] © Springer Nature Switzerland AG 2020 L. Zhu, C. Y. Li (eds.), Liquid Crystalline Polymers, Polymers and Polymeric Composites: A Reference Series, https://doi.org/10.1007/978-3-030-43350-5_58

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Abstract

Main-chain-type liquid crystalline conjugated polymers and supramolecules are discussed as hairy-rod-type polymers with rigid main chain and flexible covalently bonded side chains and as hairy-rod-type supramolecules with physically bonded side chains. Thanks to the rigid main chain these materials show both thermotropic and lyotropic liquid crystalline (LC) state which allows their facile macroscopic alignment leading to anisotropic optical and electrical properties. This chapter is a brief overview on this materials class with illustrative examples including polyfluorene, polypyridine, and polythiophene. Keywords

Conjugated polymers · Hairy-rod polymers · Hairy-rod supramolecules · Liquid Crystallinity · Alignment · Polyfluorene · Polypyridine · Polythiophene

Definition Liquid crystalline main-chain-type conjugated polymers consist of rigid main chains and flexible side chains. These materials can be discussed in terms of hairy-rod polymers and hairy-rod supramolecules and they manifest both lyotropic and thermotropic liquid LC states allowing facile macroscopic alignment.

Introduction Electrically conductive and luminescent conjugated polymers and oligomers constitute the base for organic electronics (Baeg et al. 2013; Guo et al. 2013; Wang et al. 2012). Ionic conductivity, enhanced electrochemical activity, and specific interactions towards other chemicals are achieved through additional charged groups incorporated into their structure (Jiang et al. 2009; Scherf 2011). Fig. 1 plots chemical structures of conjugated polymers discussed in this chapter. These polymers are chemically simple and serve as model compounds for generalizations while the device applications require further chemical modifications. Liquid crystals (LC) other than conjugated polymers are thoroughly discussed elsewhere in this book. Following the same principles (Akagi 2009; Bridges et al. 2017; Funahashi et al. 2008; Kuei and Gomez 2017; San Jose and Akagi 2013; Yang and Hsu 2009), main-chain conjugated polymers exhibit thermotropic nature as shown early for P1 (Tashiro et al. 1991) and lyotropic nature as shown for P2 (Wang et al. 1993) and P3 (Chen et al. 2004; Ou-Yang et al. 2005). LC state is manifested in the nanometer length scale and stems from the main-chain stiffness, while the atomic details, which are decisive for optoelectronic properties, play a secondary role. This means that so far the LC behavior is concerned, we can learn

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319

O

R R1

S

n

n

R2

P1

n

O

P2

P3

N S

n

R P4

n

R

N

n

P5

P6

SO3-Na+

S

N

S n

P7

O n

P8

S

n

P9

Fig. 1 Chemical structures of conjugated polymers mentioned in this chapter. P1: poly (3-alkylthiophene) P3nT; P2: polydiacetylene PDA; P3: poly[2-methoxy-5-(2-ethylhexyloxy)1,4-phenylenevinylene] MEH-PPV; P4: poly[9,9-bis(2-ethylhexyl)fluorene-2,7-diyl] PF2/6; P5: poly(9,9-dialkylfluorene) PFn;P6: poly(9,9-dioctylfluorene-alt-benzothiadiazole) F8BT; F7: poly (9,9-dioctylfluorene-alt-bithiophene) F8 T2; P8: poly(2,5-pyridinediyl) PPy; and P9: poly [2-(3-thienyl)ethyloxy-4-butylsulfonate] PTEBS. R stands for a linear alkyl chain

much from simple polymers and assume essentially similar trends for their chemically more complicated variants. This chapter illustrates some essential factors behind the LC nature of conjugated polymers and supramolecules especially when related in our own experience. It discusses concepts of hairy-rod polymers and hairy-rod supramolecules and compares our ideas to the works of other groups. We emphasize that LC conjugated polymers are not limited to these materials but can be prepared using various other strategies including side-chain mesogens (Liang and Lin 2009; Roudini and Foot 2016) or stacked-conjugated macromolecules (Hanack and Lang 1994; Laschat et al. 2007) or immersing-conjugated polymers into LC solutions (Fritz and Scholes 2003; Tcherniak et al. 2008; Zhu and Swager 2002), or polymerizing

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ionic liquids (Lee et al. 2010), just to mention a few. They are also discussed by K. Akagi in ▶ Chap. 12, “Liquid Crystalline Conjugated Polymers with Optoelectronic Functions.”

Conjugated Polymers as Liquid Crystalline Hairy-Rod Polymers Concept Despite the variety of chemical structures, the phase behavior of the polymers shown in Fig. 1 or similar does possess certain universality. This universality stems from the archetypical structural features of the molecules, which can be expressed via a concept of a so-called “hairy-rod molecule.” Fig. 2 plots a side view schematics of a hairy-rod molecule and head view illustrations of some possible microphases. Theory of hairy-rod polymers has been introduced in Ballauff (1986) and Stepanyan et al. (2003) and adapted to conjugated polymers in Knaapila et al. (2005b). This concept involves a rigid backbone with the length L and diameter d with L
d. The backbone is grafted to flexible side chains with N side chain segments (beads) each having a volume pffiffiffiffiv and a segment length (Kuhn length) lK. The size of the side chain coil is Rc ¼ l K N  L. Each polymer has M repeat units with the molecular weight Mu such that the number-averaged molecular weight is Mn = MuM. The Kuhn length of polymer is l HR K . The distance between grafting points is lu. side view (a)

N,n , lK

d lu

L

head view

(d)

(c)

(b)

Hex

Nem

Mem

Fig. 2 (a) A schematic of hairy-rod polymer. L and d are the length and diameter of the rod and lu the distance between grafting points. N, v, and lK are the number and volume of the side chain beads and the Kuhn length. (b) Various thermotropic and lyotropic microstructures viewed from the polymer head: Hexagonal (Hex), nematic (Nem), and “smectic” membrane (Mem) structure

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Thermotropic Behavior Limit Between LC Conjugated Polymers and LC Conjugated Oligomers Thermotropic LC conjugated polymers are crystalline becoming an ordered liquid with heating. It is important to consider the limit between LC hairy-rod polymers and LC oligomers. This consideration means the phase behavior as a function of molecular weight, Mn. Depending on Mn, the polymers manifest at least Nem and Hex phases. The phase equilibrium of the system can be understood from an analysis of the corresponding free energies. The free energy of the Nem phase is given as (Stepanyan et al. 2003) f 4π F N  k B TVcln þ k B TVcln , e ΩN

(1)

where the translational and orientational entropy are represented by the first and second terms. Here kBT is a Boltzmann factor, V the volume of the sample, c the concentration of the molecules, f the volume fraction of the backbone, and e the Euler number. The quantity ΩN describes the degree of overall (uniaxial) alignment: The smaller it is, the more aligned is the system. The free energy of the ordered Hex phase does not involve translational entropy but the interaction between ordered molecules which stems from the inhomogeneous distribution of the side chain ends starts playing an important role. For hairy rods, it reads as F lattice  k B TV

v , v0 l 2K l u

(2)

where v0 is the volume of one repeat unit of the hairy rod. From Eq. (2), the free energy of the Hex phase is F H  k B TVcln

4π v  k B TV 2 : ΩH v0 l K l u

(3)

Due to incompressibility, the concentration c = Mu/(v0Mn) is directly related to Mn. Eqs. (1) and (3) imply that the higher the Mn, the more favorable the Hex packing. The value of Mn separating Nem and Hex phases is M n  M u

l 2K l u eΩN ln : v f ΩH

(4)

To elucidate further on the temperature dependency of the phase equilibrium, one has to consider the temperature variation of the solid angles accessible for rotational degrees of freedom,ΩN, and ΩH. For both phases, the direction of each main-chain segment can fluctuate in the solid angle as (Khokhlov 1991; Semenov and Khokhlov 1988)

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Ω0 

d : l HR K

(5)

When the temperature increases, Ω0 is also increased. A first order approximation of Ω0 is Ω0 

d l HR K

ð1 þ Ct Þ,

(6)

where C is a phenomenological constant and t = T  Tg the temperature measured from the glass transition temperature Tg. The space accessible for fluctuations of the chain is  Ω i ðt Þ 

d

l HR K

ð1 þ C i t Þ

M n lu =M u lHR K ,

(7)

where i = N, H refers to “Nem” or “Hex.” This allows us to estimate the ratio ΩN ð t Þ  ΩH ð t Þ



1 þ CN t 1 þ CH t

M n lu =M u lHR K

:

(8)

The Eqs. (4) and (8) yield   M n0  t A 1 , M n

(9)

where M n0   M u

l 2K l u e ln , f v

(10)

makes a difference between low molecular weight (LMW) and high molecular  weight (HMW) polymers and where A ¼ lnðe=f Þl HR K =ðl u =ðC N  C H Þ=M n0 Þ includes all the phenomenological constants. Essentially, Eq. 10 defines the limit between hairy-rod polymers and oligomers. Figure 3 shows the phase diagram of polyfluorene P4 as a function of temperature and Mn as reported in Knaapila et al. (2005b). Solid line shows the theoretical prediction according to Eq. 9 and symbols represent experimental data. Also shown are schematics of underlying structures. Theory and experiment are quantitatively consistent when separating Nem and Hex phases. In addition, structural experiments show that the Hex phase and LMW Nem phase are actually built by bundles of three polymers. HMW Nem phase consist of separated polymer chains.

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(a)

HMW Nem

(b) 180

Nem

T (oC)

160

140

Hex

120

Hex LMW Nem

100

80

0 10 Mn*

50

100 Mn (kg/mol)

150

Fig. 3 (a) Molecular structure of 3-mers (gray) of rigid P4 (green). Side chains are omitted for clarity. (b) Phase diagram of P4 as a function of molecular weight. Solid line represents theoretically predicted phase transition between Nem and Hex phases (Eq. 9). Open and solid symbols represent experimental data from X-ray scattering and DSC. The vertical bar shows occasionally seen hexagonal traces in X-ray data. Insets illustrate organization of 3-mers in Hex and LMW Nem phases and individual hairy rods in HMW Nem phase. (Adapted and reproduced with permission from Knaapila et al. (2005b). Copyright 2005 The American Physical Society)

Lyotropic Behavior Limit Between LC Conjugated Hairy-Rod Polymers and LC Conjugated Rigid-Rod Polymers Lyotropic LC conjugated polymers are ordered liquids obtained by the addition of solvent. Fig. 4 shows a photo and polarized optical micrograph of 1% polyfluorene P5 with octyl side chains in methylcyclohexane (MCH) mixture. This is an example of a lyotropic LC conjugated polymer system. The mixture appears as a self-standing gel or a very viscous liquid whose birefringence points to the self-alignment of rod-like material. In this case, it is illustrative to consider the limit between hairy rods (with side chains) and rigid rods (without side chains). This consideration reduces to the phase behavior as a function of side chain length, N (Knaapila et al. 2008). It begins from assumptions that the polymers can be dissolved at sufficiently high temperatures forming an isotropic Iso phase and that this solution ultimately macrophase separates (demixes) on cooling. It also assumes that an intermediate structure, Mem phase,

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Fig. 4 (a) A photo of a lyotropic self-standing sample of P5 with dioctyl side chains in MCH mixture. See Knaapila and Monkman (2013) for further details. Copyright 2013 Wiley Periodicals, Inc. (b) Optical micrograph with cross-polarizers of a similar sample after 33 h ageing. (Adapted and reproduced with permission from Chen et al. (2010). Copyright 2010 The American Chemical Society)

takes place in between the solution and the demixed state. The scaling of solutionMem transition temperature, Tmem, as a function of N can be estimated by comparing the free energies of both phases. The free energy of the Iso phase contains four contributions ðsol Þ

ðsol Þ

ðsol Þ

ðsol Þ

F ðsolÞ ¼ F brush þ F RS þ F RA þ F tr ,

(11)

ðsol Þ

where the brush term F brush includes the excluded volume interaction between the alkyl side chains in a good solvent and their stretch (which has been calculated in Ref. (Subbotin et al. 2000)). Per side chain this contribution reads as ðsol Þ

F1

 ¼ kBT

νN a2 b

1=2 ,

(12)

where v is the excluded volume of a side-chain monomer in a good solvent and N and a (marked as lK in Fig. 2) the number of side chain beads and its statistical segment length. Here b is the grafting distance along the backbone (marked as lu in Fig. 2). ðsol Þ ðsol Þ The terms F RS and F RA represent the interaction of the backbone rod (“R”) with solvent (“S”) molecules and alkyl chains (“A”). The last term describes the translational free-energy of the hairy-rod polymers. The rod-solvent interaction free energy is given as (Stepanyan et al. 2003) ðsol Þ

F RS ¼ 2Ldγ RA v0 cðsolÞ , N tot

(13)

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325

where the “interfacial” energy parameter γ RS describes the interaction between backbone and solvent. Ntot is the total number of molecules in the system, ν0 the monomeric volume, and c(sol) the concentration of the solvent molecules in the vicinity of the backbone. Similarly, the rod-alkyl side-chain interaction free energy is given as ðsol Þ

F RA ¼ 2Ldγ RA v0 c2 , N tot

(14)

where  c2 is the  concentration of the alkyl monomers around the backbone so that ðsol Þ v0 c2 þ c2  1 . The last term in the Eq. (11) accounts for the translational entropy which has the Flory-type form ðsol Þ F tr

 ¼ k B TN tot

f 1f ln þ e v0 f

 2   πd L v0 NL 1f þ ln , 4 b e

(15)

where f is the volume fraction of the polymer. Thus, the Eq. 11 takes the form     F ðsolÞ L vN 1=2 γ RS γ RA f ¼ v0 c2 þ ln þ 2Ld ð 1  v0 c2 Þ þ 2 e N tot k B T b a b kBT kBT  2  1  f πd L v0 NL 1f þ , þ ln v0 f 4 b e

(16)

with c2 = (a2b2d2v)1/3 calculated based on the formulas given in Ref. (Subbotin et al. 2000). The free energy of other competing phase, the solution with membranes, can be calculated along the same lines leading to   F ðmemÞ 2NL  v 2=3 γ RS γ RA ¼ v0 c1 þ Ld ð1  v0 c1 Þ þ b abd N tot k B T kBT kBT   1  f πd 2 L v0 NL 1f þ , þ ln v0 f 4 b e

(17)

where c1 = v1/3(abd)2/3 is the concentration of the alkyl segments around a double ðmemÞ ¼ 2N ðv=ðabd ÞÞ2=3 for is the planar layered sheet of backbones. Moreover, F 1 brush free energy per chain (Milner et al. 1988). A critical temperature Tmem, below which a transition from isotropic solution to Mem phase is expected to occur, is obtained by equating Eqs. (16) and (17) 1=2 pffiffiffiffi N  ðb=LÞlnðf =eÞ 1 2ðv=abd Þ2=3 N  ðv=a2 bÞ :  ¼ bd ½γ RS þ ðγ RA  γ RS Þv0 ð2c2  c1 Þ k B T mem

(18)

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Here we assume that ν, γ RA, and γ RS do not depend on the temperature. Although generally incorrect, these assumptions do not change the scaling behavior of Tmem. The results are valid for relatively long side chains, at least so long that the distance between their grafting points is smaller than their Flory radius RF = a(v/a3)1/5N3/5, i.e., N > N = (b/a)5/3(v/a3)1/3. For shorter side chains, one can expect a picture where isotropic liquid demixes with decreasing temperature (Ballauff 1986). Following similar ideas, the free energy of the macroscopically demixed state is estimated as F ðdemixÞ 3NL  v 2 γ ¼ þ Ld RA , 2b abd N tot k B T kBT

(19)

where it is assumed that a layered lamellar structure in the polymer-rich phase is ðlamÞ formed and F 1 ¼ 3=2k B TN ðv0 =ðabd ÞÞ2 is the elastic energy per side chain (see Stepanyan et al. 2003 for details). Accordingly, the demixing temperature, TIN is

1 k B T IN 

 v 1=2 pffiffiffiffi b  f 1  f πd 2 L v NL 1  f  3  ν 2 0 þ N  ln þ N 2 ln 2 bad a b L e v0 f 4 b e : ¼ bd ½γ RS þ ðγ RA  γ RS Þð1  v0 c2 Þ (20)

Figure 5 plots a theoretical phase diagram of a lyotropic hairy-rod mixture and the experimental phase diagram of P5 in MCH mixture as a function of N as reported in Knaapila et al. (2008). The phase transition temperatures are predicted by Eqs. (18) and (20). In both cases, Mem phase dominates lower temperatures and shorter side chains and Iso phase higher temperatures and longer side chains. Similarly, Nem phase is observed for polymers approaching rigid rods. Iso phase is not attainable for sufficiently dense solutions of rigid rods. These results are in a good qualitative agreement with the experimental phase diagram. Interestingly, the free energy of Mem phase, Eq. (17), scales as the free energy of demixing, Eq. (19). This implies that the membranes are likely thermodynamically unstable, which is consistent with the experimental observations (Knaapila et al. 2008).

Alignment Limit Between LC Conjugated Polymers and LC Conjugated Oligomers LC conjugated polymers can be aligned using methods that require their LC state such as annealing directly on a pretreated substrate (Nothofer et al. 2000) or the friction transfer method (Misaki et al. 2004), where the pretreated substrate is covered by a second layer which does not eliminate the orienting ability of the substrate. These methods involve mechanically treated “rubbed” substrates while

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327

(a) TIN*(N) T

isotropic solution

nematic & solvent coexistence

Tmem*(N)

membranes nematic & solvent coexistence or metastable membranes N*

N

360

(b) isotropic solution

T (K)

320

300

nematic & solvent coexistence

340

metastable membranes

280

260

Tmem *(N)

6

7

8

9

10

N Fig. 5 (a) Theoretical phase diagram of lyotropic hairy-rod polymer and (b) experimental phase diagram of P5 in dense MCH mixture as a function of side chain beads, N. (Adapted and reproduced with permission from Knaapila et al. (2008). Copyright 2008 The American Physical Society)

soft lithography allows another route for the substrate preparation (Zheng et al. 2007). LC conjugated polymers can also be aligned by methods that do not require the LC state per se, such as epitaxial crystallization (Brinkmann 2007) or floating transfer (Dauendorffer et al. 2012).

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While polymers in composites are usually aligned to improve their mechanical performance, conjugated polymers are aligned to improve their optoelectronic performance. Fig. 6 shows the schematics of alignment procedure and polarized photoluminescence (PL) data of P6 as reported by Zheng et al. (2007). PL becomes polarized in the alignment procedure along the alignment direction. This is already seen in images taken with a polarizer parallel and perpendicular to the alignment direction. Fig. 7 plots the transfer characteristics and charge carrier mobilities as a function of temperature for aligned P7 thin-film transistors (TFTs). The mobility values are much higher when parallel than perpendicular to the alignment direction, and the values of non-aligned device lie between these two extremes. It is again illustrative to consider the limit between LC conjugated polymers and oligomers (Knaapila et al. 2005a). First, we clarify what the previously discussed Ω means in the terms of the order parameter s (for further details see the other chapters of this book or de Gennes and Prost (1998)). The molecule, whose rigid backbone is defined by the vector c, is free to rotate in the solid angle Ω so that it can take any orientation with angle θ with the director between 0 and θ0, where Ω = 4π(1  cos θ0) (Fig. 8). In the ordered state, the angle distribution function f(θ) can be approximated as f ðθÞ ¼

c~, if 0

0 < θ < θ0 otherwise

or

π  θ0 < θ < π

(21)

The constant c~ is found from the normalization of f and equals c~ ¼ 0:5=ð1  cos θ0 Þ. The order parameter s and the angle Ω are, in turn

Fig. 6 (a) Schematics of aligning LC conjugated polymers using nanoconfinement. (b) PL spectra of aligned P6 parallel (red circles) and perpendicular to the alignment direction (black triangles) alongside corresponding PL images. (Adapted and reproduced with permission from Zheng et al. (2007). Copyright 2007 The American Chemical Society)

Liquid Crystalline Conjugated Polymers

a

329

300 Vsd =–60V 200

Id [nA]

II ×

100

0

⊥ –50

–30

–10

–20

0

Vg [V]

b

10–2

5

10–3

4

II ⊥

3

10–4

3.5

Fig. 8 Relation between the molecular axis c and the alignment direction z. (Reproduced with permission from Knaapila et al. (2005a). Copyright 2005 The American Chemical Society)

–40

μ ll / μ⊥

Fig. 7 (a) Transfer characteristics (Vsd = 60 V) of aligned P7 TFTs with channels parallel (upward triangles) and perpendicular (downward triangles) to the alignment direction. Circles show corresponding data for a nonaligned device. (b) Temperature dependence of the linear (open symbols) and saturated (closed symbols) mobility (μ) and mobility anisotropy (μk/μ⊥) of devices with channels parallel and perpendicular to the alignment direction. (Reproduced with permission from Sirringhaus et al. (2000). Copyright 2000 The American Institute of Physics)

μ [cm2/Vs]

11

4

4.5 1000/T [K–1]

5

5.5

z

q0

c

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   3 cos2 θ  1 1 Ω Ω ¼ 1 s ¼ dθf ðθÞ 2 : 2 2 4π 4π 0 pffiffiffiffiffiffiffiffiffiffiffiffiffi Ω 3  1 þ 8s ¼ : 4π 2 ðπ

(22)

(23)

Here s = 0 and s = 1 refer to the completely isotropic state and perfect alignment with Ω = 4π and Ω = 0. The order parameter is related to the dichroic ratio for absorption, which is defined as



R ¼ Ek =E⊥ ,

(24)

where Ek and E⊥ are the maximum values of the absorbance for light polarized parallel and perpendicular to the alignment direction z and potentially parallel to the molecular c axis (Fig. 8). R describes the anisotropy of the absorption process. The transition probability is maximized when the transition moment of the molecule lies parallel to the electric vector of the light. This moment is assumed to be parallel to the c axis. R is related to the order parameter as s = (R  1)/(R + 2). Therefore, for large R, Eq. (23) takes the form  Ω 2  þ O R2 : 4π R

(25)

This result shows that if Ω increases exponentially with Mn, as Eq. (7), then R decreases exponentially. Fig. 8 plots dichroic ratio R for P4 as a function of molecular weight. When the polymer is denoted as HMW material, it follows theoretical prediction whereby R decreases exponentially with increasing Mn. In contrast, when the polymer is denoted as LMW material, the orientation becomes better and thus R increases with increasing Mn. This shows a major difference in the behavior of hairy-rod oligomers and polymers (Fig. 9).

Conjugated Polymers as Liquid Crystalline Hairy-Rod Supramolecules Thermotropic Behavior LC conjugated supramolecules can be conceptually understood in terms of hairy-rod supramolecules. Fig. 10 plots a side view schematics of a hairy-rod supramolecule and head view illustrations of some possible microphases in the solid state. Like in the case of hairy-rod polymers, the backbone and side chains are both chemically and geometrically different and tend to form rod-rich and coil-rich microdomains. The driving force behind this microphase separation – the unfavorable interaction between the stiff backbone and the flexible side chains – can be described by the surface tension γ which is proportional to χδ, where χ is the Flory-Huggins parameter and δ is the width

Liquid Crystalline Conjugated Polymers

Hex

35

331

100

30

R

Dichroic Ratio in Absorption

40

Nem

11

10

25 20 1 10

15 10

50 Mn (kg/mol)

90

5 0 0 Mn* 20

40

60 80 100 Mn (kg/mol)

120

140

Fig. 9 Dichroic ratio in absorption of aligned P4 as a function of molecular weight. Inset shows the relation between the molecular axis c and the alignment direction z. (Reproduced with permission from Knaapila et al. (2005a). Copyright 2005 The American Chemical Society)

of the interpenetration region between the pure rod and pure coil phase. As discussed in Subbotin et al. (2003), the formation of microphases can be understood from an analysis of the free energies of the different phases. In general, the free energy F is written as F γ F el ffi Sþ , T T T

(26)

where S is the interface area describing unfavorable contacts between rods and side chains and Fel the stretching energy describing stretching of side chains. The former term dominates for short and the latter for long side chains, but both also depend on the geometry of the microphases. The free energy can be calculated from the number of side chain beads, N, and the ratio between the volumes of coil and backbone sections between two consecutive branching sites, κ = v/(πd2b/4). Minimization of free energy against the rod fraction f corresponds to various microstructures. Moreover, a Nem structure is expected for very short side chains and an isotropic high temperature structure for very long side chains. Unlike covalently bound hairy-rod polymers, hairy-rod supramolecules can macrophase separate to rod-rich and coil-rich phases or all the way to pure constituents because of the limited mixing entropy of rod-like polymers (Flory 1984). Microphase separation occurs only if the association energy between rods and coils e is high

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(a)

N, a, ν

–ε d b (b)

Sq (Hex1)

Obl (Hex2)

Lam (Lam)

Fig. 10 (a) A schematic of hairy-rod supramolecules. N, a, and v are the number, length, and the volume of side chain beads. b is the distance between active sites on a rigid rod and d is the rod diameter. e is the association energy between side chains and backbone. (b) Various microstructures viewed from the polymer head: Square, Sq, elliptical oblique, Obl, and lamellar, Lam, structure. (Adapted and reproduced with permission from Knaapila et al. (2003). Copyright 2003 The American Chemical Society)

enough. This means that the real phase diagram contains domains for both micro- and macrophase separation and should be plotted as a function of T and f for a given e. An example of conjugated hairy-rod supramolecule demonstrating both microand macrophase separation is given in Knaapila et al. (2003). This involves complexation of polypyridine P8 by methane sulfonic acid (MSA) and amphiphile octyl gallate (OG). Fig. 11 shows polarized optical micrographs of this material above and below the order-disorder transition (ODT) into an Iso phase when the degree of complexation is 0.75. LC domains are yellow, while Iso phase becomes black. This system corresponds to small e and short side chains which together mean microand macrophase separated phase regimes and dominance of Lam microstructure. Fig. 12 plots corresponding theoretical and experimental phase diagrams. The experimental case follows the theoretical prediction but contain also Hex2 phase (or an Obl phase) and square Sq phase that is not theoretically expected for this particular system.

Lyotropic Behavior Solvent is another factor in the phase behavior of conjugated hairy-rod supramolecules. A prime example that illustrates not only lyotropic but also thermotropic phase

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Fig. 11 Optical micrographs with cross-polarizers of P8 complexed with MSA and OG. (y = 0.75) above (a) and below (b) an order-disorder transition. (Adapted and reproduced with permission from Knaapila et al. (2003). Copyright 2003 The American Chemical Society)

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behavior is given by Osuji and coworkers (Zhang et al. 2010). This involves stoichiometric complexation of aqueous polythiophene P9 by cetrimonium bromide (CTAB). Figure 13a shows a photo of this mixture after centrifugation showing, from top to bottom, fluid-like, gel-like, and slurry phases. This phase separation is attributed to the wide polydispersity with the highest molecular weight fraction populating the bottom phase. The attraction through charge transfer between SO3 and N(CH3)3+ groups is strong (and thus ise high). Hence, the separation occurs between the solvent and complex and not between the backbone and side chains. Fig. 13b shows a photo of the middle layer demonstrating gel-like nature with high viscosity. The structure of gel depends on the supramolecule fraction and temperature. As an example, Fig. 14 shows two-dimensional X-ray scattering pattern and polarized optical micrograph of shear aligned 70% gel. In this polymer fraction, the material shows Hex phase with LC nature as indicated by strong birefringence. Fig. 15 shows schematics of this example and the overall phase diagram as a function of temperature and polymer fraction (or concentration). The thermotropic behavior from crystalline structures to the LC phase is a special case with 100% polymer concentration.

Challenge Although the chemical synthesis of main-chain-type conjugated polymers has gone through tremendous advancements, much of their LC phase behavior can be understood in terms of hairy-rod polymers and hairy-rod supramolecules, whereby most of the chemical details can be effectively ignored. In this work, polyfluorenes have proved to be one successful model compound for both thermotropic and lyotropic behavior. We believe that this knowledge provides a starting point for various expansions.

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Iso

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Fig. 12 (a) Theoretical phase diagram of hairy-rod supramolecule and (b) experimental phase diagram of P8 complexed with MSA and OG. (Adapted and reproduced with permission from Knaapila et al. (2003). Copyright 2003 The American Chemical Society)

Theoretical work could be moved towards kinetics and nonequilibrium states. For example, the alignment by annealing on the pretreated substrates is currently continued until the alignment is saturated and the saturated values are compared to the theoretical prediction. Similarly, theoretically monitored experiments of lyotropic systems are performed immediately after sample preparation. The theory of ageing and phase instabilities remains incomplete. So far the quantitative work is focused on materials bulk. Both theoretical and experimental work should have stronger emphasis on thin films and multilayers where the polymers are influenced by confinement effects and by interfaces of other

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SO3-

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Fig. 13 Photos of the aqueous P9 complexed with CTAB. (a) Three layers (fluid, gel, and solid) after centrifugation. (b) Gel from the middle layer. (Adapted and reproduced with permission from Zhang et al. (2010). Copyright 2010 The American Chemical Society)

Fig. 14 (a) X-ray scattering pattern and (b) optical micrographs with cross-polarizers of shearaligned gel of P9 complexed with CTAB. White arrows indicate the shearing direction. (Reproduced with permission from Zhang et al. (2010). Copyright 2010 The American Chemical Society)

material layers. We also suggest that the field should be expanded towards alignment by external electric and magnetic fields. While this area is well developed for low molecular weight LCs, much more could be done with conjugated polymers as already demonstrated for conjugated polymers in photovoltaic devices (Chen et al. 2014; Zhao et al. 2011).

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Fig. 15 (a) A schematic of hairy-rod supramolecule and its Hex microstructure. (b) Phase diagram for aqueous P9 complexed with the CTAB amphiphile. Cry1, Cry2, and LC stand for two different crystalline forms and a thermotropic LC phase in the solid state (100% polymer concentration). (Adapted and reproduced with permission from Zhang et al. (2010). Copyright 2010 The American Chemical Society)

Materials for organic electronics involve conjugated stiff segments and their processing steps include many times solutions and annealing. This means that the LC nature, whether intentional or not, becomes an inherent part of their materials science and engineering. Already for this reason, we believe that deepening understanding of LC phase behavior of conjugated polymers and supramolecules remains important for years to come.

Cross-References ▶ Liquid Crystalline Conjugated Polymers with Optoelectronic Functions

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Liquid Crystalline Conjugated Polymers with Optoelectronic Functions

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Kazuo Akagi

Contents Introduction . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . Generation of Monodomain Structures of Mono-substituted LCPAs . . . . . . . . . . . . . . . . . . . . . . . . . . Structures and Liquid Crystallinity of the Mono-LCPAs . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . Macroscopic Alignment and Linearly Polarized Luminescence of Di-substituted LC Polyacetylene Derivatives . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . Synthesis and Structures of the Di-LCPAs . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . Absorption and Photoluminescence of the Di-LCPAs . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . Liquid Crystallinity of the Di-LCPAs . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . Linearly Polarized Luminescence of the Di-LCPAs . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . Summary . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . Cross-References . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . References . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . .

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Abstract

We review the recent progress in advanced functionalities of liquid crystalline polyacetylenes (LCPAs). We focus on properties and functionalities that include electrical anisotropy and linearly polarized luminescence (LPL). First, the synthesis of mono-substituted LCPAs using Fe-based Ziegler–Natta catalyst is briefly reviewed. The mono-LCPAs exhibit enantiotropic SA phases resulting from spontaneous orientation of the LC side chain. Iodine doping of monoLCPA cast films followed by macroscopic alignment of the main chain accompanied by the side chain orientation using an external magnetic force of 0.7–1.0 T enhanced the electrical conductivity by two orders of magnitude to 10 6 S/cm and gave rise to a notable electrical anisotropy. Second, the synthesis of disubstituted LCPAs using metathesis catalyst that exhibit enantiotropically K. Akagi (*) Research Organization of Science and Technology, Ritsumeikan University, Kusatsu, Japan e-mail: [email protected] © Springer Nature Switzerland AG 2020 L. Zhu, C. Y. Li (eds.), Liquid Crystalline Polymers, Polymers and Polymeric Composites: A Reference Series, https://doi.org/10.1007/978-3-030-43350-5_46

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thermotropic or lyotropic liquid crystallinity is discussed. LC phases of di-LCPAs are assigned through observation of polarized optical microscope (POM) and differential scanning calorimeter (DSC) and measurements of X-ray diffraction (XRD). Using schematic energy levels of ground and low-lying excited states of non-, mono-, and di-PAs, the origin of the emission of substituted PAs is elucidated, and fluorescent trends including emission color are investigated. It is found that the macroscopically aligned films of the di-LCPAs emit LPL by virtue of the functionalities associated with liquid crystallinity and fluorescence. The aligned structures of the di-LCPAs are characterized in terms of main chain and side chain type alignments through XRD measurements of the macroscopically aligned polymer films. The mechanism of the LPL of the di-LCPAs with respect to the polymer structure, alignment type, and emission color is elucidated. Keywords

Liquid crystallinity · Conjugated polymers · Substituted polyacetylenes · Linear polarization · Luminescence

Introduction Polymers are generally considered insulators. However, it has been demonstrated that conjugated polymers can become electrically conductive. Among these conjugated polymers, polyacetylene (PA) has attracted much attention due to the discovery of its metallic conductivity in the doped form. When doped with iodine, pristine PA film has a metallic luster with a black surface and a high electrical conductivity on the order of 104–105 S/cm (Naarrmann and Theophilou 1987; Akagi et al. 1989; Tsukamoto et al. 1990). This discovery accelerated research on conjugated polymers and led to developments in polymeric light-emitting diodes (PLEDs), plastic electronics, polymer battery cells, polymer photovoltaics, and other novel technologies (Skotheim and Reynolds 2007; Perepichka and Perepichka 2009; Mullen et al. 2014). However, pristine PA film is insoluble in organic solvents and quickly loses its electrical conductivity when exposed to atmospheric conditions. Introduction of an alkyl substituent into the polymer main chain increases the solubility in organic solvents depending on the length of the alkyl chain. However, the electrical conductivity of the substituted PA is significantly lower than that of non-substituted PA. This phenomenon is due to decreased coplanarity of the main chain, which arises from steric repulsions between the substituents, a higher ionization potential and lower electron affinity. The main chain in the substituted PA remains randomly oriented, which suppresses the electrical conductivity of the polymer. The addition of a liquid crystalline (LC) moiety to the polymer main chain makes the polymer soluble in organic solvents and facilitates alignment by the spontaneous orientation of the LC group. For example, mono-substituted liquid crystalline PA (mono-LCPA), which is prepared by introducing a LC moiety into the side chain

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of the PA, exhibits liquid crystallinity and stability in air (Moigne et al. 1992; Oh et al. 1993; Vicentini et al. 1994; Koltzenburg et al. 1998; Choi et al. 2000; Ting et al. 2002; Liu et al. 2009b). It is well known that polymers with LC moieties spontaneously align to form multidomains by virtue of the orientation of the LC moieties. Further linear arrangement of these polymers can be achieved by applying an external force, such as shear stress and electric or magnetic fields, to construct a monodomain structure on a macroscopic level. The LC polymers in a monodomain structure exhibit excellent anisotropic properties including electrical conductivity and linearly polarized luminescence (LPL). Figure 1 shows the formation of multidomains, monodomains, and helical assemblies through spontaneous orientation as well as linear and helical arrangement of LC conjugated polymers. It has been reported that mono-LCPA derivatives exhibit excellent macroscopic alignment under an applied magnetic field. Macroscopic alignment of the main chain led to both an enhancement of two orders in electrical conductivity and an electrical anisotropy accompanied by the magnetically forced orientation of the LC side chain. Developments in the macroscopic alignment and anisotropic electrical conductivity

Fig. 1 Schematic representation of the spontaneous orientation, linear, and helical arrangement of LC conjugated polymers

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of mono-LCPA as well as representative results of this research have been investigated in detail (Akagi et al 1995; Akagi 2007, 2009). Compared to aromatic conjugated polymers, such as poly(para-phenylene) (PPP) and polythiophene (PT), mono-substituted PA (mono-PA) is generally considered non-emissive due to the coulombic interactions between π-electrons (Shukla and Mazumdar 1999). Figure 2 depicts the energy levels of the ground and low-lying excited states of non-, mono-, and di-substituted PAs and the correlations between the corresponding energy states. In Fig. 2, 1Ag, 2Ag, and 1Bu denote the intrachain excitons. These excitons are conserved in mono- and di-substituted PA derivatives, but their symmetries are different from non-substituted PA. Therefore, for this series of PA derivatives, the optical transitions involving the Ag and Bu pair of states are symmetry allowed. However, the optical transitions involving two Ag states are symmetry forbidden because the transition dipole moment in the direction of the main chain is strongest and its irreducible representation is Bu. In trans-PA, which has C2h symmetry (in terms of unit cell), the transition between the ground state 1Ag and the excited state 2Ag is electronically forbidden. However, according to the symmetry-based selection rule, the transition between 1Ag and the excited state 1Bu is electronically allowed. For non- and mono-substituted PAs, the 1Bu state decays into the 2Ag state, but the transition from 2Ag to 1Ag is electronically forbidden. Therefore, non- and mono-substituted PAs are generally non-emissive. When another bulky group is substituted into the mono-PA to form di-substituted PA (di-PA), the steric effect on the polyene main chain changes the relative position

Fig. 2 Energy levels of ground and low-lying excited states of non-, mono-, and di-substituted polyacetylenes and correlations between the corresponding energy states (San Jose 2011)

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of the two excited states such that the 1Bu excited state is lower than the 2Ag state. This shift enables a radiative transition from the 1Bu state to the 1Ag ground state (Ghosh et al. 2000; Hidayat et al. 2001; Shukla et al. 2001). Therefore, fluorescent di-PA can be synthesized by introducing a second side chain into the mono-PA main chain. The addition of LC moieties to di-PA gives di-substituted LC PA (di-LCPA) whose mesophase has either thermotropic or lyotropic liquid crystallinity. By combining the ability of LC-substituted polymers to be aligned macroscopically and the fluorescent functionality of di-PA, di-LCPA is able to exhibit LPL. Here, we discuss the development of di-LCPA derivatives that exhibit LPL functionality and the LPL behavior in terms of the polymer structure, alignment, and emission color. In this entry, we survey the development of LC conjugated polymers by focusing on advanced electrical and optical properties of di-LCPA that exhibit anisotropic polarized functionalities.

Generation of Monodomain Structures of Mono-substituted LCPAs Mono-PA has a significantly lower electrical conductivity than non-substituted PA due to the random orientation of its main chains and other factors. To increase the electrical conductivity of mono-PA, an LC moiety was introduced to the side chain of mono-PA for LC macroscopic alignment. The first report of mono-LCPA (PA1), which contains the bulky cholesteryl group, was reported by Le Moigne et al. (1992) (Scheme 1). Their X-ray diffraction (XRD) results show that the mono-LCPA forms a smectic A (SA) phase with a layer period corresponding to the length of the side chain in its extended conformation. Akagi et al. have also synthesized a series of mono-LCPAs (PA3) with a mesogenic moiety of phenylcyclohexyl group (Scheme 2) (Oh et al. 1993). The PA3 derivatives present the typical fan-shaped texture of SA phase, and their phase transition is found to be enantiotropic. Using the phenyl benzoate moiety as the LC mesogen, Vicentini et al. also obtained PA2, which exhibits a SA phase that is stable over a large temperature range (Scheme 1) (Vicentini et al. 1994). These pioneering studies opened the avenue to mono-LCPA research. Mono-LCPA is easily aligned through the spontaneous orientation of the LC group. In addition, mono-LCPA is macroscopically aligned by an external perturbation, such as shear stress and electric or magnetic fields, resulting in a monodomain structure for the LC phase that can be constructed on a macroscopic level. Therefore, the polymer has a higher electrical conductivity compared to the random orientation case. In addition, one can simultaneously control the molecular orientation and the electrical conductivity of the polymers with an external force.

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Scheme 1 Mono-substituted LCPA derivatives with side chain type liquid crystallinity

Structures and Liquid Crystallinity of the Mono-LCPAs A series of mono-LCPA derivatives including PA3 and PA4 (Scheme 2) have been synthesized (Oh et al. 1993). The LC substituent on the monomers was composed of a phenylcyclohexyl (PCH) moiety for PA3 or a biphenyl (BP) one for PA4 as the mesogenic core, a methylene chain linked with an ether-type oxygen atom ( (CH2)3O) as the spacer, and an alkyl chain (–CmH2m + 1, m = 5, 8) as the terminal group. The polymers exhibited thermal stability and thermotropic liquid crystallinity with an enantiotropic nature. PA3 and PA4 exhibited fan-shaped textures and uniaxial conoscopic patterns under a polarizing optical microscope (POM) that are characteristic of the SA phase (Dierking 2003). For PA3, the transition temperature from the SA to the isotropic phase increased as the terminal alkyl chain length increased, which indicates that the liquid crystal state is stabilized better by the longer alkyl chains used as the terminal moiety of the LC group. It was demonstrated that mono-LCPA derivatives were uniaxially aligned from the shear stress or magnetically forced alignment of the LC side chain to yield a monodomain structure. The electrical conductivity of the mono-LCPA derivatives was enhanced to 10 6 S/cm and gave rise to a notable electrical anisotropy. Additional functional aligned mono-LCPA derivatives have also been developed (Scheme 3). Tang et al. synthesized PA5 containing phenyl benzoate mesogens with octyl spacers and cyano or methoxy tails that exhibit well-ordered parallel bands after the application of shear force (Kong and Tang 1998). Akagi et al. developed

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Scheme 2 Structures of conducting mono-substituted LCPA derivatives

PA6 and PA7 with ferroelectric liquid crystalline (FLC) functional moieties (Goto et al. 2004; Suda and Akagi 2008; Akagi 2009). The mono-LCPAs bearing FLC moieties exhibit a characteristic chiral smectic C phase (SC), which has the ability to quickly respond to an electric field. Hsu et al. extended the concept of macroscopic alignment to chiral saccharide-containing PA8 (Ho and Hsu 2009). The helical twist morphological structure of PA8 can be aligned on a polyimide rubbed glass substrate to form two-dimensional ordered helical patterns. In addition, Okoshi et al. synthesized optically active main chain type mono-LCPAs having a poly (phenylacetylene) (PPA) main chain structure and D- or L-alanine moieties (PA9) (Okoshi et al. 2005). The rigid PPA main chain, which has bulky phenyl group and the intramolecular hydrogen bonding between the amide groups in the neighboring side chains, can exhibit cholesteric phases in concentrated solutions. The development of these mono-LCPAs led us to extend the use of LC functionality to disubstituted PAs where the fluorescence of the di-PAs can be combined with liquid crystallinity to gain novel functionalities, such as linearly and circularly polarized luminescence.

Macroscopic Alignment and Linearly Polarized Luminescence of Di-substituted LC Polyacetylene Derivatives Over the past few years, various functional di-substituted polyacetylene derivatives were developed. Masuda et al. developed gas permeable di-PA films (Masuda et al. 1988; Sakaguchi et al. 2008) and investigated di-PAs with photoelectronic properties (Shiotsuki et al. 2011). Hsu et al. reported the electroluminescence (EL) of novel

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Scheme 3 Structures of functional mono-substituted LCPA derivatives

fluorene containing di-PAs for applications in light-emitting diodes (LEDs) and organic light-emitting diode (OLEDs) (Huang et al. 2006). Tang et al. reported diPAs with liquid crystallinity, highly efficient photoluminescence, and aggregationinduced emission (Liu et al. 2009). Furthermore, Chen et al. and Tang et al. developed mesogen jacketed liquid crystalline polyacetylene (MJLCPAs), with mesogen cores laterally linked to the side chain of polyacetylene through very short spacers showing unique columnar LC phases (Peng et al. 2010; Yu et al. 2013). Earlier reports of LPL in di-LCPAs have been reported by Kwak and Fujiki et al. (Kwak et al. 2007, 2008). PA10, which has a stiff poly(diphenylacetylene) (PDPA) main chain structure and an alkyl tail for lyotropic liquid crystallinity (Scheme 4), was aligned using shear stress and electrospinning. Recently, we developed di-LCPA with novel structures where the LC mesogens are either substituted directly into the main chain or connected by flexible spacers (San Jose et al. 2011). In this section, we describe the macroscopic alignment of the di-LCPAs and the mechanism of their LPL behavior with respect to their main chain and side chain structures.

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Scheme 4 Structures of PDPA derivatives showing linearly polarized luminescence

Synthesis and Structures of the Di-LCPAs We have synthesized various di-LCPA derivatives, including two groups with structural variations on the alkyl side chain, the LC mesogen core, and the terminal groups (Scheme 5). The Group 1 and Group 2 di-LCPAs have side chain type liquid crystallinity, where the LC mesogens are located on their side chains. The Group 1 di-LCPAs have the terphenyl LC mesogen core directly attached to the polyene main chain, while the Group 2 di-LCPAs have the LC groups attached via flexible spacers. PA11 has an n-hexyl group and a para-terphenyl mesogen core linked with an n-octadecyloxy moiety. PA12 has a phenyl group and a para-terphenyl mesogen core linked with an n-pentyl moiety. The di-LCPAs with the second group have the mesogen cores attached to the polyene main chain via flexible spacers. PA13 has a pentyl-PCH group attached with an n-nonyloxy flexible spacer and an n-pentylphenyl group. PA14 has two phenyl groups directly attached to the polyene main chain with the second phenyl having a PCH group attached via an n-dodecyloxy flexible spacer. PA15 has a similar structure to PA14, with the phenyl group having a PCH moiety attached via a n-dodecyloxy flexible spacer. On the other side of the PA15 main chain is a PCH group with an n-nonyloxy flexible spacer. The synthesis of the monomers and polymers is described in the report (San Jose et al. 2011). The monomers of the di-LCPA derivatives were polymerized via a metathesis reaction using tantalum (V) pentachloride (TaCl5) as a catalyst and tetra-nbutyltin (n-Bu4Sn) as a co-catalyst. The polymerization reaction was conducted at 80  C under an argon atmosphere using toluene as a solvent. The polymerization proceeded for 24 h followed by washing of the polymer in methanol (MeOH) under constant stirring for 24 h. The precipitated polymers were then washed in acetone for 24 h and dried under vacuum. It has been extensively studied and reported by Masuda et al. that sterically demanding monomers such as ortho-substituted phenyl acetylenes polymerize

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Scheme 5 Structures of the di-substituted LCPA derivatives showing linearly polarized luminescence

with metathesis catalysts to yield high molecular weight di-substituted PAs (Shiotsuki et al. 2011). They have reported the synthesis of freestanding PDPA gas permeable films, using TaCl5 as catalyst, with very high number average molecular weights (Mn = over 105 Da). The Group 1 and Group 2 di-LCPAs were synthesized using TaCl5/n-Bu4Sn catalyst to yield polymers with Mn and the polydispersity (Mw/Mn) values ranging from 19,000 to 210,000 Da and 1.7 to 4.5, respectively.

Absorption and Photoluminescence of the Di-LCPAs We examined the UV-vis and photoluminescence of the di-LCPAs in chloroform (CHCl3) solution and cast film. PA13, which has a poly(alkylphenylacetylene) (PAPA) main chain structure, shows a main chain absorption band at 320 nm and a PCH mesogen absorption band at 280 nm (Fig. 3a). PA11 and PA15, having also

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Fig. 3 UV-vis and PL spectra of (a) PA13 and (b) PA14 in solution and in cast film, insets: PL in solution (left) and in cast film (right) (San Jose 2011)

PAPA main chain structures, show similar absorption spectra. On the other hand, PA14, which has a PDPA main chain structure, exhibits three absorption peaks at 280, 380, and 430 nm (Fig. 3b) corresponding to the absorption of the PCH moiety, the trans-stilbene structure, and the polyene main chain, respectively. PA12, which has a PDPA main chain structure, also exhibited two absorbance peaks corresponding to the polyene main chain and the trans-stilbene moiety. The PL spectra of PA11, PA13 (Fig. 3a), and PA15 show blue-colored emission, with their maximum emission (Emmax) being around 480 nm. The PA14, with a PDPA main chain structure, has green fluorescence with an Emmax around 500 nm (Fig. 3b). PA12 also exhibits green luminescence at 520 nm. Previously, we elucidated the origin of emission of substituted PAs and investigated the fluorescent trends of di-LCPA. In the earlier work, we observed that di-LCPAs with PAPA and PDPA structures generally exhibit blue and green emissions, respectively. The green emission color of the latter is associated with the PDPA structure in which two bulky phenyl rings are directly attached to the polyene main chain. Specifically, the two bulky side chains contribute to lower the 1Bu excited state and hence to decrease the band gap (Fig. 2) (Shukla and Mazumdar 1999; Ghosh et al. 2000), resulting in a bathochromic shift of the emission bands of PA12 and PA14. Generally, the PAs show small emission shifts between the solution and cast film, which is especially advantageous for applications of di-PAs as fluorescent materials in EL devices.

Liquid Crystallinity of the Di-LCPAs The di-LCPAs exhibited thermotropic and lyotropic liquid crystallinity. The Group 1 and Group 2 di-LCPAs show side chain type liquid crystallinity where their LC phases are summarized in Table 1. POM analysis revealed that PA11 has a N-LC phase with a Schlieren texture (Fig. 4a) and no isotropic phase. The temperature ranges of the N-LC phase were

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Table 1 Thermotropic and lyotropic liquid crystallinity of the di-substituted LCPAs Polymers PA11

Liquid crystal phase Thermotropic

PA12 PA13

Lyotropic Thermotropic

PA14

Thermotropic

PA15

Thermotropic

G 125 N over G 100 N 250 N (10–15 wt% in toluene) G 75 SA over G 70 SA 210 G 90 N over G 85 N 250 G 55 SB 120 I G 55 SB 110 I

G glassy state, N nematic phase, SA smectic A phase, SB smectic B phase, I isotropic phase

125–250  C and 100–250  C in heating and cooling processes, respectively, and thermal decomposition occurred at temperatures over 250  C. The LC phase of PA11 was also investigated using XRD. The XRD patterns reveal a single broad diffraction peak in the wide angle region (Fig. 4b), typical of an N-LC phase (Shibaev and Lam 1994). The broad peak at 19.6 in 2θ is 4.5 Å, and it is assigned to the distance between mesogenic groups. The two PCH mesogen cores attached via flexible spacers on PA15 result in a smectic B (SB) phase (Fig. 4c) with temperature ranges of 55–120  C and 55–110  C in heating and cooling processes, respectively. The XRD patterns of PA15 have two sharp diffraction peaks at 5.2 and 19.4 in 2θ (Fig. 4d), corresponding to distances of 34.4 and 4.6 Å, respectively. The former represents the smectic interlayer distance, while the latter corresponds to the distance between mesogenic groups. The sharp signal in the wide angle peak of the XRD profile is typical in a hexagonal SB phase due to the regularity of the hexagonal packing arrangement. Among the Group 2 di-LCPAs, PA15 has a lower transition temperature (55  C) from the glassy to LC phase than that of PA13 (75  C) and has an isotropic phase. This is attributed to the two PCH LC groups in the side chains of PA15. Conversely, PA14 has a higher transition temperature (90  C) from the glassy to LC phase than that of PA13 (75  C), because of the main chain stiffness brought about by the directly attached phenyl and phenylene groups in the side chains. Meanwhile PA12 has no thermotropic LC phase; however, it exhibits lyotropic liquid crystallinity. To achieve a lyotropic LC phase, a stiff polymer structure with high molecular weight and a good solubility toward a solvent under an appropriate concentration at room temperature are needed (Shibaev and Lam 1994). The main chain of PA12 has a stiff polymer structure because its side chains are composed of phenyl and terphenyl moieties directly attached to the main chain forming a stilbene fragment enforced by π-conjugation. A polymer exhibiting this elevated level of stiffness should be infusible upon heating, resulting in an LC with no thermotropic behavior. Instead, by virtue of the alkyl terminal group substituted at their side chains, PA12 is soluble in organic solvents such as a toluene and shows lyotropic LC behavior. Note that the other Group1 and Group 2 di-LCPAs (PA11, PA13-PA15) are also soluble in organic solvents, but they show no lyotropic LC. This finding is

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Fig. 4 (a) POM image of PA11 at 120  C in cooling process. Inset shows a Schlieren texture of the nematic LC (N-LC) phase. (b) The XRD pattern of PA11 shows a broad reflection at 4.5 Å (19.6 in 2θ). Inset shows the Laue pattern of PA11. (c) POM image of PA15 at 110  C in cooling process. (d) The XRD pattern of PA15 shows sharp reflections at 34.4 Å (5.2 in 2θ) and at 4.6 Å (19.4 in 2θ). Inset shows the Laue pattern of PA15. (e) POM image of PA12 showing a lyotropic N-LC phase at 10 wt% toluene solution. (f) XRD pattern of PA12 showing a broad reflection at 5.0 Å (17.6 in 2θ). Inset shows the Laue pattern of PA12 (San Jose 2011)

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probably due to their higher solubility, which might depress the formation of spontaneously aligned self-assemblies that is a prerequisite to the formation of domains of lyotropic LC. In other words, polymers with high solubility are freely dissolved in a solvent with randomly oriented conformation and form neither a regularly stacked nor an associated structure. This situation is far from what is necessary for the formation of the lyotropic LC. In the case of the PDPA main chain structure of PA12, the stilbene fragment might be suitable for the formation of an interchain π-electron overlapped association through van der Waals interactions, which enables the polymers to exhibit the lyotropic LC. PA12 demonstrates lyotropic liquid crystallinity at a critical concentration range from 10 to 15 wt% in toluene. Analogous to thermotropic LCs which show an isotropic phase above a critical temperature, PA12 becomes an isotropic liquid below 10 wt% concentration in toluene. At the isotropic phase, the polymer main chains (which act as LC mesogens) are dispersed randomly without any ordering. At the lyotropic LC concentration range, there are weak π-interactions between the aromatic stilbene structure and toluene. In such a situation, the toluene enhances the attraction between polymer main chains for spontaneous self-assembly to occur and also provide enough fluidity in the system for liquid crystallinity to take place. The lyotropic LC phase of PA12 was prepared from a 10 wt% solution using toluene as a solvent. Figure 4e shows the POM image of the polymer solution, depicting an optical texture characteristic of the N-LC phase. The XRD analysis of PA12 shows a single broad diffraction peak at 17.6 in 2θ. The peak corresponds to 5.0 Å, which is assigned to the distance between the mesogenic groups (Fig. 4f).

Linearly Polarized Luminescence of the Di-LCPAs The macroscopic alignment of the Group 1 and Group 2 di-LCPA films was achieved using rubbing technique. Polymers on quartz substrates were heated to the LC temperature region and rubbed with a glass rod along the long axis of the quartz substrate before cooling to room temperature. In the case of PA12, lyotropic LC solution in 10–15 wt% toluene was prepared, then applied to a quartz substrate, and then rubbed with a glass rod along the long axis. The PL intensity was measured as a function of the polarizer angle with respect to rubbing direction. The dichroic ratio (DR) of the di-LCPAs is summarized in Table 2 and is defined as the ratio of parallel to perpendicular polarized PL intensity (DR = I// / I⊥) or vice versa (DR = I⊥ / I//). Three patterns of alignment and emission behavior were observed on the aligned PA films (see Fig. 6 later). The alignment behavior of the PA films can be related to the structure of the polymer main chain and the LC moieties. In the di-LCPA derivatives, the main chain and the side chain can both act as LC mesogens. The interactions between the main chain and the LC moieties determine the alignment behavior of the PA films. On the other hand, the emission behavior of the PA films can be related to the diPAs exciton confinement (localization) - deconfinement (delocalization) mechanism

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Table 2 Linearly dichroic ratios of the di-substituted LCPAs Polymersa Group 1 Group 2

PA11 PA12 PA13 PA14 PA15

Dichroic ratio 1.6 2.2 1.2 2.4 1.6

Aligned films of PA11 and PA13-PA15 were prepared in thermotropic LC states and that of PA12 was prepared in lyotropic LC state

a

(Shukla and Mazumdar 1999; Ghosh et al. 2000; Hidayat et al. 2001; Shukla et al. 2001). Excitons are electron-hole pairs that are formed during photoexcitation. For the di-LCPAs, the excitons are initially confined (localized) to the polyene main chain because of steric effects between the bulky side chains (Fig. 5a). They then are deconfined (delocalized) to the phenyl moieties where they radiatively recombine (Fig. 5b). The structures of the polyene main chain and the phenyl moieties of the diPAs determine the polymer emission behavior. In Case 1, as represented by PA11, we observe that PA11 has a DR of 1.6 and a higher PL intensity perpendicular to the rubbing direction (Fig. 6a). The reflection of the para-terphenyl moiety was observed at the meridian parallel to the rubbing direction in the XRD profile of the aligned PA11 film (Fig. 6b). These results imply that the para-terphenyl moiety and hence the polyene main chain are aligned perpendicular and parallel to the rubbing direction, respectively (Fig. 6c). For PA11, the excitons are moved to the para-terphenylvinyl moiety through deconfinement; thus the dominant emission of PA11 is from the para-terphenylvinyl moiety with the LPL perpendicular to the rubbing direction. The DR of PA13 is 1.2, with a higher PL intensity perpendicular to the rubbing direction (Fig. 6d). The XRD pattern of aligned PA13 film shows a meridional reflection at the small angle area of 5.7 in 2θ corresponding to the main chain interlayer distance of 31 Å (n = 2). An equatorial reflection at 18.0 in 2θ was also observed (Fig. 6e). This reflection can be assigned as the distance of 4.9 Å between LC moieties aligned parallel to the rubbing direction. Given the PA13 LPL spectra and XRD profile, we propose an alignment structure where the polyene main chain and the LC moieties are perpendicular and parallel to the rubbing direction, respectively (Fig. 6f). In PA13, the PCH LC moiety aligns parallel to the rubbing direction, and thus it exhibits side chain type LC alignment. For PA13, the excitons are moved to the phenylvinyl moieties along the main chain. As a consequence, the dominant emission of PA13 is from the phenylvinyl moieties. The Case 2 di-LCPAs, PA13 and PA15, show the same side chain type LC alignment and emission behavior. In Case 3, as represented by PA14, we observe that PA14 has a DR of 2.4 and a higher PL intensity parallel to the rubbing direction (Fig. 6g). The XRD pattern of aligned PA14 film shows an equatorial reflection at the 17.8 in 2θ (d = 5.0 Å) corresponding to the distance between LC moieties (Fig. 6h). The LPL data and the

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Fig. 5 Exciton (a) confinement (localization) and (b) deconfinement (delocalization) behavior in di-substituted LCPA (San Jose 2011)

XRD results imply that the PCH LC moieties and hence the polyene main chain are aligned parallel and perpendicular to the rubbing direction, respectively (Fig. 6i). In PA14, the PCH LC moiety aligns parallel to the rubbing direction, and thus it exhibits side chain type LC alignment. The stilbene moieties aligned parallel to the rubbing direction are the dominant emission sources; thus PA14 exhibits a higher PL intensity parallel to the rubbing direction. Figure 7 shows an image of the LPL of PA14 under polarizing plates. PA12 and PA14 exhibit side chain type LC alignment and dominant emission from the stilbene moiety. The PAs show LPLs with fairly low DR values that range from 1.2 to 2.4. These values are mainly due to the situation characteristic of the di-LCPAs where not only the LC side chain but also the conjugated polyene main chain directly linked with phenyl or phenylene moiety behave as mesogens to the external force. Since the LC side chain and the main chain are almost orthogonal in the polymer structure, the externally forced alignments of the side and main chains are substantially cancelled each other out, thus giving a low degree of alignment. From the above results, we can deduce that when one designs emissive polyacetylene derivatives with linearly dichroic nature, it is essential for the polymers to have phenylene moieties that are directly or indirectly attached to the main chain and also to introduce LC moieties in one side or both side chains. However, these polymers inevitably encounter the competitive orientation or even partial

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Fig. 6 (a) Linearly dichroic PL of aligned PA11, (b) XRD pattern of aligned PA11 shows a broad peak at 4.5 Å (19.6 in 2θ) corresponding to the distance between the terphenyl LC moieties in the side chains, inset: XRD Laue pattern of aligned PA11 film, and (c) orientation diagram of aligned PA11. (d) Linearly dichroic PL of aligned PA13, (e) XRD pattern of aligned PA13 shows peaks at 31 Å (n = 2, 5.7 in 2θ) and 4.9 Å (18 in 2θ) corresponding to the interlayer distance of the main chains and the distance between the PCH LC moieties in the side chains, respectively, and (f) orientation diagram of aligned PA13. (g) Linearly dichroic PL of aligned PA14, (h) XRD pattern of aligned PA14 shows a peak at 5.0 Å (17.8 in 2θ) related to the distance between the PCH LC moieties in the side chains, and (i) orientation diagram of aligned PA14. Rectangles with orange broken lines in (c, f, and i) indicate dominant emission moieties in the polymers (San Jose 2011)

cancellation in alignment between the two mesogenic cores of the rigid polyene main chain and the LC side chain that are almost orthogonal to each other. Among Group 1 and Group 2 di-LCPAs, the polymers bearing the double mesogenic core in side chains (PA12, having the stilbene and terphenyl mesogens; and PA14, having stilbene and PCH mesogens) are the most promising for the dichroic emissive PAs,

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Fig. 7 Photograph showing the LPL of PA14, where two polarizers are arranged parallel (bright) and perpendicular (dark) to the aligned direction. The white arrow indicates the alignment direction (San Jose 2011)

because the double mesogenic core is highly responsive to the external perturbation than the single mesogenic one, giving an alignment with substantially notable linear dichroism.

Summary In this entry, a broad overview of the recent developments in LC polyacetylene derivatives and their linearly polarized functionalities was discussed. The structures and liquid crystallinity of mono-LCPA derivatives were briefly reviewed. The synthesis of di-LCPA derivatives using metathesis catalyst that exhibit enantiotropically thermotropic or lyotropic liquid crystallinity was discussed. The origin of the emission of substituted PAs was elucidated, and fluorescent trends including emission color were investigated. It was found that the macroscopically aligned films of the di-LCPAs emit LPL by virtue of the functionalities associated with liquid crystallinity and fluorescence. The aligned structures of the di-LCPAs were characterized in terms of main chain and side chain type alignments through XRD measurements of the macroscopically aligned polymer films. The mechanism of the LPL of the di-LCPAs with respect to the polymer structure, alignment type, and emission color was elucidated. Over the past few years, research focused on LC polyacetylenes has been overlooked in favor of LC aromatic conjugated polymers. However, recent developments with the LC polyacetylenes described in this entry may encourage further development of di-LCPAs with advanced LC functionalities. The synergistic incorporation of functional di-LCPAs with other advanced polymer functionalities, such as ferroelectric liquid crystallinity as well as temperature- and photo-switching, might lead to the emergence of novel next-generation LC conjugated polymers.

Cross-References ▶ Anisotropic Liquid Crystal Networks from Reactive Mesogens ▶ Characterizations of Nanocomposites of Liquid Crystalline Polymers ▶ Columnar Phase-Forming Polymers ▶ Liquid Crystalline Conjugated Polymers

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▶ Liquid Crystalline Polymers Derived from Disc-Shaped Molecules ▶ Mesogen-Jacketed Liquid Crystalline Polymers: Molecular Design and Synthesis ▶ New Stimuli-Response Liquid Crystalline Polymer Architectures ▶ Photoactive Liquid Crystalline Polymers ▶ Photoalignment of Liquid Crystal Molecules Using Fluorine-Containing Polyimides ▶ Photodeformable Liquid Crystalline Polymers (LCPs) ▶ Photoresponsive Liquid Crystalline Polymers ▶ Structure and Assembly of Liquid Crystalline Block Copolymers Acknowledgments The author is grateful to Dr. Benedict A. San Jose and Dr. Satoshi Matsushita (Department of Polymer Chemistry, Kyoto University) for their valuable contributions in the syntheses of liquid crystalline conjugated polymers. This work was supported by a Grant-in-Aid for Science Research (A) (No. 25246002) and (No. 25620098) and (No. 15K13706) from the Ministry of Education, Culture, Sports, Science and Technology, Japan.

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Photodeformable Liquid Crystalline Polymers (LCPs)

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Lang Qin, Wei Gu, and Yanlei Yu

Contents Definition . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . Introduction . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . Photodeformation Based on Photochemical Phase Transition . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . Photodeformation Driven by Visible and NIR Light . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . Photodeformation Based on Photothermal Effect . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . Soft Actuators . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . Macroscaled Actuators . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . Microscaled Actuators . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . Summary . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . References . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . .

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Abstract

Collecting and amplifying the nanoscopic molecular motions into macroscopic deformation are the basic properties of crosslinked liquid crystalline polymers (CLCPs), which can even directly transfer input light energy into mechanical work when combined with photochromophores, thus fascinating many scientists. This article reviews the macroscopic and microscopic deformation of photoresponsive CLCPs based on photochemical phase transition and photothermal effect. In addition, we highlight some new methods to trigger the deformation driven by visible and infrared light instead of ultraviolet one, such as chemical modification of azobenzene and addition of upconversion materials.

L. Qin · W. Gu · Y. Yu (*) Department of Materials Science and State Key Laboratory of Molecular Engineering of Polymers, Fudan University, Shanghai, People’s Republic of China e-mail: [email protected]; [email protected]; [email protected] © Springer Nature Switzerland AG 2020 L. Zhu, C. Y. Li (eds.), Liquid Crystalline Polymers, Polymers and Polymeric Composites: A Reference Series, https://doi.org/10.1007/978-3-030-43350-5_52

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Keywords

Crosslinked liquid crystalline polymers · Photoresponsive · Deformation · Photochemical phase transition · Photothermal effect · Actuators

Definition Crosslinked liquid crystalline polymers (CLCPs) are a type of promising smart materials that possess both the ordered alignment of liquid crystals (LCs) and the elasticity of polymer networks. The anisotropic photodeformation of CLCPs takes place when mesogens experience order to disorder change in response to light. Therefore, many actuators are fabricated from CLCPs and have potential applications in artificial muscles, micro-optomechanical systems, optics, and energy-harvesting fields.

Introduction Smart materials have drawn wide attention recently due to their unique properties and are potential for applications in artificial muscles and soft actuators, biomedical systems, etc. Among these smart materials, CLCPs appear to be especially attractive because of their superior characteristics combining the order of LCs and the excellent mechanical properties arising from polymer networks. Because the mesogens tend to show alignment in CLCPs coupled with the elastic property of the polymer networks, the CLCPs as three-dimensional networks are able to undergo controllable and reversible deformation due to the alignment change induced by external stimuli. Compared to other stimulusdriven method, such as heat, pressure, pH variations, electric field, and magnetic field, light is a particularly ideal stimulus, since it is a clean energy and can be precisely and conveniently manipulated in terms of wavelength, intensity, and polarization direction. The incorporation of photoresponsive chromophores into CLCPs can provide photoresponsiveness and induce a reduction in LC alignment upon exposure to UV light as a result of photochemical reaction (Ikeda et al. 2007; Ube and Ikeda 2014; White and Broer 2015). Therefore, photodeformable CLCPs can convert light energy into mechanical actuation directly and definitely merit further investigation. In this chapter, we mainly describe photoinduced deformation observed in CLCPs. The mechanisms of deformation based on photochemical phase transition and photothermal effect are also included. Our goal is to summarize the developments of photoinduced behavior of CLCPs and provide an introduction on their potential applications as light-driven devices as well as recent progress in this field.

Photodeformation Based on Photochemical Phase Transition Cooperative motion of molecules in LC phases may be most advantageous in changing the molecular alignment by external stimuli. The alignment of the majority of LC molecules will be changed if the alignment of a small portion of LC molecules

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is changed in response to an external stimulus. This phenomenon illustrates that LC molecules only require a small amount of energy to change the alignment: the energy needed to induce an alignment change of only 1 mol% of the LC molecules is enough to bring about the alignment change of the whole system. In other words, a huge amplification is possible in LC systems. When a small amount of a photochromic molecule is added into LCs and the resulting guest/host mixture is irradiated to cause photochemical reactions of the photochromic guest molecules, LC-isotropic phase transition of the mixtures can be induced isothermally. Ikeda et al. reported the first explicit example of nematic-isotropic phase transition induced by trans-cis photoisomerization of a nematic LC with an azobenzene guest molecule dispersed in it (Tazuke et al. 1987). Azobenzene is a well-known chromophore which has two configurations. It undergoes trans to cis photoisomerization upon exposure to UV irradiation, and irradiation with visible light leads to cis to trans back-isomerization process. Therefore, azobenzene is the most frequently used photochromic moiety in photoresponsive polymers. The rodlike trans form of the azobenzenes stabilizes the structure of LC phase, whereas its bent cis isomer tends to destabilize the phase structure of mixture (Fig. 1a). As a consequence of two different conformations, the LC-isotropic phase transition temperature (Tc) of the mixture with the cis form (Tcc) is much lower than that with the trans form (Tct). If the temperature of the sample (T) is between Tct and Tcc and the sample is irradiated to cause trans-cis photoisomerization of the azobenzene guest molecules, Tc decreases because of the increasement of the cis form. When Tc becomes lower than the irradiation temperature T, LC-isotropic phase transition of the sample is induced. The sample reverts to the initial LC phase through cis-trans back-isomerization due to reversible photochromic reactions. Thus, phase transitions of LC systems can be induced isothermally and reversibly by photochemical reactions of photoresponsive guest molecules (Fig. 1b) (Ikeda 2003). Finkelmann et al. reported pioneering work on photodeformation of a monodomain nematic CLCP, which had a polysiloxane main chain and azobenzene chromophores at crosslinks. The CLCP film generated a contraction by 20% upon irradiation with UV light to give rise to the trans-cis isomerization of the azobenzene moieties (Fig. 2) (Finkelmann et al. 2001). It is necessary to take photomechanical effects into consideration: the subtle variation in nematic order upon trans-cis isomerization causes a significant uniaxial deformation of the LCs along the director axis when the LC molecules are strongly associated by covalent crosslinking to form a three-dimensional polymer network. The contracted elastomer thermally returned to the original state due to the cis-trans back-isomerization after stopping irradiation. Keller and coworkers synthesized oriented monodomain nematic side-on CLCPs containing azobenzenes by photopolymerization with a near-infrared photoinitiator (Li et al. 2003). The photopolymerization was performed with aligned azobenzene monomers in conventional LC cells. The obtained thin films were found to show fast (less than 1 min) photochemical contraction of up to 18% upon exposure to UV light and a slow thermal recovering in the dark. In order to observe a strong effect of the isomerization, high concentrations of the photoresponsive molecules are required. This will lead to a high optical density of

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Fig. 1 (a) Schematic illustration of reversible LC-isotropic photochemical phase transition. (b) Phase diagrams of the photochemical phase transition of azobenzene/LC systems (N nematic, I isotropic) (Ikeda 2003)

photoresponsive molecule-containing film. The light absorption intensity within the film is no longer constant, and the extent of isomerization varies throughout the sample. Due to the different degrees of elongation or contraction, the internal stress can lead to strong bending deformation (Fig. 3). Furthermore, the azobenzene moieties are preferentially aligned along the rubbing direction of the alignment layers, and thus the decrease in alignment order of azobenzene moieties is produced along this direction, which contributes to the anisotropic bending behavior (Ikeda et al. 2007). Ikeda and coworkers were the first to report photoinduced bending behavior of macroscopic CLCPs containing azobenzene (Ikeda et al. 2003). Compared to the two-dimensional contraction or expansion, the bending mode, three-dimensional movement, can be advantageous for a variety of real manipulation applications.

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Fig. 2 (a) Chemical structures of the monomers and crosslinker. (b) Photoinduced contraction of CLCP. ⎕ = 313 K, △ = 308 K, ○ = 303 K, * = 298 K. (Inset) Recovery of the contracted CLCP at 298 K after irradiation was switched off (Finkelmann et al. 2001)

The monodomain CLCP film bent toward the irradiation source along the rubbing direction and reverted to the initial flat state upon exposure to visible light. This bending and unbending behavior was reversible and controlled by simply altering the wavelength of the incident light. Furthermore, when the film was rotated by 90 , the bending behavior was again observed along the rubbing direction. These results illustrate that the bending is anisotropically induced and occurs only along the rubbing direction.

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Fig. 3 Plausible mechanism of the photoinduced bending of CLCP film. (a) Bending toward the light source. (b) Bending away from the light source. (Ikeda et al. 2007)

Ikeda and coworkers prepared the films by thermal polymerization of a liquid crystal monomer and a diacrylate crosslinker. By means of selective absorption of linearly polarized UV light in polydomain CLCP films, they succeeded in controlling the direction of photoinduced bending so that a single polydomain CLCP film was found to be bent repeatedly and precisely along any chosen direction (Fig. 4a) (Yu et al. 2003). The film bent toward the irradiation source in a direction parallel to the polarization of the light and completely reverted to its initial flat state upon exposure to visible light with a wavelength longer than 540 nm. The polydomain CLCP film consists of many micro-sized domains of azobenzene moieties aligned in one direction in each domain. Although macroscopically the direction of alignment is random (Fig. 4b), upon exposure of the film to linearly polarized light, the selective absorption of light in a specific direction leads to a contraction in specific domains where the azobenzene moieties are aligned along the direction of light polarization. It is well-known that human skeletal muscles are composed of many bundles of fibers and their crucial function is to convert chemical energy into mechanical work.

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Fig. 4 (a) Chemical structures of liquid crystal monomer and diacrylate crosslinker. (b) Precise control of the bending direction of a film by linearly polarized light: photographs of the polydomain film in different directions in response to irradiation by linearly polarized light at different angles of polarization (white arrows) at λ = 366 nm; the bent films are flatted by irradiation with visible light at λ > 540 nm. (c) Schematic illustration of the plausible bending mechanism (Yu et al. 2003)

Ikeda et al. prepared CLCP fibers containing an azobenzene moiety by two-step reactions (Yoshino et al. 2010). It was found that the CLCP fibers exhibited a Tg of around 60  C and showed a high order of mesogens along the fiber axis. When the CLCP fiber was irradiated with UV light, the CLCP fiber bent toward the actinic light source along the fiber axis. The bent fiber recovered to the initial state upon exposure to visible light. The photoinduced bending and unbending of the CLCP fiber was reversible simply by changing the wavelength of the actinic light, similar to that of CLCP films. Furthermore, a three-dimensional control of bending direction in the CLCP fibers was carried out with the experimental setup shown in Fig. 5b. Since the shape of the CLCP fiber was approximately cylindrical, the direction of the bending could be controlled by changing the irradiation direction of the actinic light. The generated stress upon contraction of the natural surface length reached 210 kPa, which is similar to the stress in human muscles (around 300 kPa). In order to achieve the orientation in the CLCP films, generally, an aligned polyimide layer with parallel grooves generated by mechanical rubbing along one direction was often used to orient the LC molecules. Lately, by using highly aligned carbon nanotube (CNT) sheets, a new and general method to prepare photodeformable CLCP/CNT nanocomposite films was developed (Wang et al. 2012).

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Fig. 5 (a) Chemical structures of the monomers and crosslinker. (b) Schematic illustration of experimental setup. (c) Photographs of the CLCP fiber that exhibits photoinduced bending and unbending behavior upon irradiation with UV light (100 mW cm2) and visible light (120 mW cm2). The inset of each photograph is a schematic illustration of the state of the fiber. The size of the fiber is 30 mm  20 μm (Yoshino et al. 2010)

The CLCP/CNT composite film exhibited a rapid and reversible deformation under alternate irradiation by UV and visible light (Fig. 6). This actuation is derived from the structure change in the composite film, which results from the photoisomerization of the azobenzene moieties. Compared to the CLCPs prepared by

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Fig. 6 (a) Chemical structures of two monomers and crosslinker. (b) Preparation of an oriented CLCP/CNT nanocomposite film. (c) Photographs of a CLCP/CNT composite film during one bending and unbending cycle after alternate irradiation by UV light at 365 nm (100 mW cm2) and visible light at 530 nm (35 mW cm2), respectively (Wang et al. 2012)

the conventional mechanical rubbing method, the introduction of aligned CNTs remarkably improved mechanical strength and high electrical conductivity of the CLCP film. Besides the bending behavior, the coiling movement of CLCPs in response to light was also reported. Broer and coworkers prepared CLCP films with a densely crosslinked, twisted configuration of azobenzene units (Harris et al. 2005). Although the networks were stiff and glassy at room temperature, the films showed large amplitude coiling motion as well as bending motion upon exposure to UV light, which was based on the configuration of twisted LC alignment of 90 . Recently, Iamsaard and coworkers reported complex motion of springlike CLCP materials (Iamsaard et al. 2014). Nature provides a valuable source of inspiration for many fields of research. Based on the general concept of plantlike helical deformations, such as spasmoneme springs, seed pod opening, and tendril coiling, they designed, synthesized, and studied the versatile actuation modes of photoresponsive CLCP springs (Fig. 7). A small amount of chiral dopant S-811 was added into the mixture to induce a left-handed twist. The resultant orientation of LC director smoothly changed by 90 from bottom to top surface (Fig. 7a). The direction in which the ribbon was cut determined the pitch, the handedness of the helical shapes, and their photoresponsive behaviors. Under irradiation with light, left-handed spiral ribbons doped with S-811 decreased in their macroscopic pitch, and the corresponding right-handed ribbons showed an increase in macroscopic pitch. Remarkably, it was also possible to observe inversion of the helical

Fig. 7 (a) Molecular organization in the twist cell (top view) and the angular offset φ, which characterizes the angle at which the ribbon is cut. The orientation of the molecules at midplane is shown with a double-headed arrow. The cutting direction, which is also the long axis of the ribbon, is represented by a dotted line. The elongated rods represent molecules (left) and the twist-nematic molecular orientation through the thickness of the film (side view) (right). (b) Schematic

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sense from right-handed to left-handed. As a result, these springs displayed complex motion, including winding, unwinding, and helix inversion which depended on the handedness of the director twist and on their angular offset φ (Fig. 7c). The ribbons always deformed to accommodate the preferred distortion along the main axis of the ribbon, and this preferred distortion was determined by the orientation of the molecules, which was, in turn, determined by the cutting direction. The mixed-helicity springs comprising two opposite-handed helices displayed unwinding and winding motion simultaneously under UV irradiation, which successfully mimicked the movements of plant tendrils (Fig. 7d).

Photodeformation Driven by Visible and NIR Light To develop applications of light-driven organic actuators in possible biological systems, low-energy light instead of UV light would be a more suitable stimulating source because low-energy light penetrates deeper into tissues and causes less damage to biosamples. Moreover, as the stimulating source, UV light is not environment-friendly and causes harm to our health, limiting its practical applications. Furthermore, sunlight is the origin of all the energy resources that can be endlessly supplied, and visible light is harmless and more abundant in sunlight. Thus, it would be useful to develop the CLCPs with photochromic molecules that undergo a photoinduced deformation in response to visible light, especially sunlight. Yu et al. first reported visible light-induced bending and unbending of azotolanecontaining CLCPs, whose deformation even occurred upon exposure to sunlight (Yin et al. 2009; Cheng et al. 2010b). Compared with 366 nm absorption of usual azobenzene moieties, the maximum absorption of the azotolane groups shifts toward a long wavelength region at 385 nm, resulting in a decrease in the energy level difference between the π and π* orbital of the tolane groups. Irradiated with shortwavelength visible light at 436 nm, the film bent toward the irradiation direction of the actinic light due to the trans-cis photoisomerization of azotolane and reverted to the initial state after irradiation with visible light at 577 nm. The azotolane CLCP film also underwent photoinduced bending and unbending behavior by means of manipulating the wavelength of sunlight through a lens and glass filters as shown in Fig. 8. This kind of sunlight-responsive film is of great importance in development ä Fig. 7 (continued) illustration showing the direction in which the ribbons are cut. (c) Spiral ribbons irradiated for 2 min with UV light (λ = 365 nm) display isochoric winding, unwinding, and helix inversion (φ was defined as the angle between the orientation of the molecules at midplane and the cutting direction; R, right-handed; L, left-handed). (d) A coiled tendril of the wild cucumber plant (left) and a polymer spring that displays a cucumber tendril-like shape, composed of two oppositely handed helices (middle and right). On irradiation the right-handed helix unwinds, and the lefthanded helix winds (Iamsaard et al. 2014)

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Fig. 8 (a) Chemical structures of the monomer and crosslinker. (a) Experimental setup. (b) Photoinduced bending and unbending behavior of azotolane CLCP film in sunlight through a lens and glass filters. The sunlight at >430 nm and at >570 nm was acquired by using different filters. (Yin et al. 2009)

and utilization of solar energy because it converts solar energy into mechanical energy directly. Furthermore, Yu et al. gave the first example to incorporate upconversion nanophosphors (UCNPs) NaYF4:Yb,Tm into the azotolane-containing CLCP film and succeeded in generating fast bending of the resulting composite film upon irradiation with continuous-wave (CW) NIR light at 980 nm (Fig. 9) (Wu et al. 2011). Here, upconversion luminescence (UCL) of the nanophosphors not only induces trans-cis photoisomerization of the azo groups but also leads to alignment change of the mesogens. Under excitation with a CW 980 nm laser, the as-prepared UCNPs show blue emission, and the main UCL emission peaks at 450 nm and 475 nm, as shown in Fig. 9a, overlap the absorption band of the azotolane CLCP film (between 320 nm and 550 nm) perfectly; thus the UCL light emitted by UCNPs triggers the trans-cis photoisomerization of the azotolane moieties. This kind of novel photodeformable CLCP system is promising for biological applications, since NIR light penetrates deeper into tissues and has less damage to biosamples. Lately, Yu et al. achieved a red light-controllable composite film driven by low-power excited UCL based on triplet-triplet annihilation (TTA) (Jiang et al. 2013). This TTA-based UCL process shows several advantages over the lanthanide upconversion techniques, such as higher quantum efficiency, large absorption efficiency, and low excitation power density. When PtTPBP and BDPPA were incorporated into a

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Fig. 9 (a) UCL emission spectrum (blue line) of a colloidal CHCl3 solution of UCNPs (1 mg mL1) excited with a 980 nm CW laser (power = 600 mW, power density = 15 W cm2) and the UV-vis absorption spectrum (black line) of azotolane CLCP film. The inset shows a photograph of UCL from the UCNPs in CHCl3. (b) Schematic illustration of the mechanism of CW NIR light-induced deformation of the azotolane CLCP/UCNP composite film and photographic frames of the composite film bending in response to the NIR light at CW 980 nm and being flattened after removing the light source (Wu et al. 2011)

soft polyurethane film and then assembled with an azotolane-containing CLCP film, a soft material system was achieved (Fig. 10). Upon excitation of 635 nm laser, the PtTPBP&BDPPA-containing polyurethane film acts as an antenna to trap the 635 nm light and upconvert it into the blue TTA-UCL emission; then the TTA-UCL is absorbed

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Fig. 10 (a) Chemical structures of the sensitizer PtTPBP (left) and the annihilator BDPPA (right). (b) UCL emission spectrum (blue line) of toluene solution of PtTPBP&BDPPA (λex = 635 nm, power density = 200 mW cm2) and the UV-vis absorption spectrum (red line) of the azotolane CLCP film. (c) Schematic illustration of the preparation of the assembly film composed of azotolane CLCP film and PtTPBP&BDPPA-containing polyurethane film. (d) Photographs of the as-prepared assembly film bending toward the light source along the alignment direction of the mesogens in response to the 635 nm laser with the power density of 200 mW cm2 (thickness of each layer in the assembly film: 15 μm of upconverting film and 27 μm of CLCP film). (e) Schematic illustration demonstrating plausible mechanism for the photoinduced deformation of the as-prepared assembly film (Jiang et al. 2013)

by the azotolane moieties in the CLCP film via the emission-reabsorption process, which induces the trans-cis photoisomerization of the azotolane moieties and the subsequent alignment change of the mesogens, thus contributing to the photoinduced bending of the azotolane CLCP film toward the light source. Moreover, to our most interest, the assembly film still bent toward the light source even though a piece of pork with the thickness of 3 mm is put between the light source and assembly film, which demonstrates potential biological applications using this novel red light-controllable soft actuator. This work not only provides a novel photomanipulated soft actuation material

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system based on the TTA-UCL technology but also introduces a new technological application of the TTA-based upconversion system in photonic devices.

Photodeformation Based on Photothermal Effect If a CLCP film is heated from the LC phase (typically nematic) to isotropic state, the LC order decreases, and the mesogens become disordered when the temperature exceeds the phase transition temperature. With this phase transition, the CLCP films in general exhibit contraction along the alignment direction of the mesogens, while the CLCP films revert to their original shape (expansion) if the temperature is lowered below the phase transition temperature (Fig. 11). In 1975, de Gennes et al. first proposed the concept of the CLCPs as artificial muscles by taking advantage of their substantial uniaxial contraction in the direction of the director axis (De Gennes 1975). He also theoretically predicted the possibility of a large deformation of CLCPs induced by the phase transition. Finkelmann et al. prepared CLCP film with polysiloxane main chain and LC side chain, which showed a contraction of 26% when heated to isotropic state due to the change in molecular alignment (Kupfer and Finkelmann 1991). This superior property improved the ability of CLCPs to function as artificial muscles. Since then, a number of studies have been done on artificial muscle-like CLCPs (Ohm et al. 2010; White and Broer 2015). The development of thermo-driven actuators has been restricted in practical applications because of the difficulty to implement contactless and precise control. To this end, many efforts have been made through absorptive heating with either optical or magnetic stimuli. Therefore, the carbon nanoparticles, magnetic

Fig. 11 Schematic illustration of thermal-induced anisotropic contraction and expansion of the CLCPs

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Fig. 12 (a) Setup for the preparation of patterned alignment cells and procedure for making a patterned polymer LC. (b) Actuation behavior of an azimuthal CLCP film upon heating with an IR lamp. (c) Actuation behavior of a radial CLCP film. The arrows along the radius and the azimuth indicated the direction of deformation. (de Haan et al. 2012)

nanoparticles or nanorods, carbon nanotubes, and gold nanoparticles were added to the CLCPs, which offers the possibility to induce thermal phase changes locally in their environment, since the absorption of photons by carbon nanotubes or gold nanoparticles is radiated as heat and heat transfer triggers thermomechanical effects in the CLCPs. As for magnetic nanoparticles, they are known to transfer energy from electromagnetic irradiation into heat due to relaxation processes. Schenning and Broer et al. used photoalignment technique to align polymerizable LCs and prepare freestanding CLCP films with complex order (de Haan et al. 2012). Setup for the preparation of patterned alignment cells is shown in Fig. 12a. The CLCP films with azimuthal and radial alignments were successfully prepared using the LC mixture comprising an IR-absorbing dye (Lumogen 788 IR). The films deformed into conical and anti-cone shape when heated by absorbing IR light. In the case of an azimuthal alignment pattern, a reduction of the LC order upon temperature increase led to compression along the azimuthal direction and an expansion along the radial direction. Consequently, a conical shape was observed as shown in Fig. 12b. The opposite deformations took place in the case of a radial alignment pattern, resulting in an anti-cone shape (Fig. 12c). Similarly, Chen et al. prepared wavelength-selective, IR light-driven bilayer hinges composed of one active CLCP composite layer with IR-active fillers and one passive silicone layer. IR-active fillers consisted of single-walled carbon nanotubes (SWCNTs) and near-IR dyes which generated heat by absorption of light (Kohlmeyer and Chen 2013). Bilayer hinges could show fast, reversible bending with a large strain originating from the bulk CLCP nematic-isotropic phase transition. The bilayer films were fabricated into active origami structure and inchworm

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Fig. 13 (a) Scheme of an origami structure. (b) Reversible folding and unfolding of the origami structure in response to continuous-wave (CW) NIR light (11.0 mW mm2). (c) Reversible closing and opening of a Venus flytrap-inspired gripper in response to CW NIR light (28.2 mW mm2). (d) Scheme of an inchworm walker device. (e) Scheme of a ratcheted wood substrate. (f) The inchworm walker crawling up the wood substrate at a 50 incline in response to on and off cycles of CW NIR light (28.2 mW mm2) (Kohlmeyer and Chen 2013)

walker when assembled with polycarbonate (PC) films. The foldable origami structure inspired by a Venus flytrap opened and closed repeatedly in response to IR light (Fig. 13c). The inchworm walker could further crawl up the wood substrate at a 50 incline upon IR light irradiation cycles (Fig. 13e). Moreover, the gripper composed of PC films and bilayer hinges was able to pick up and place various objects. As mentioned above, the CLCP materials are capable of transferring optical or thermal energy into mechanical energy. Hence they have promising applications in artificial muscles and microelectromechanical systems (MEMs). From an energy point of view, it is also fascinating to harvest solar energy and transfer it into electricity. Jiang et al. demonstrated artificial heliotropism for solar cells, utilizing polyurethane fiber-network/SWCNT/CLCP actuators that could be directly driven by the sunlight (Li et al. 2012a). Figure 14a shows the concept of the artificial heliotropism. The solar cells are installed on a platform that is connected to the actuators. SWNTs absorb and transform the incoming sunlight into thermal energy, so that the actuators are in a contracted state toward the direction of the incident light

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Fig. 14 (a) Three-dimensional (3D) schematic of the heliotropic behavior. (b, c) The actuator facing the sun contracts, tilting the solar cell toward the sunlight. White light source intensity: 100 mW cm2

at any time instant. While other actuators, which are not exposed to the sunlight, remain in the relaxed state. Consequently, the platform holding the solar cells is driven by the contracted actuators and self-adaptively tilts toward the sunlight. The devices were capable of full-range artificial heliotropism with 60 of range in altitude angle and 180 of range in azimuth angle, and the resultant output photocurrent of solar cells increased significantly.

Soft Actuators An actuator is an energy transducer that can convert input energies of a variety of forms into mechanical quantities such as displacement, strain, velocity, and stress (Ikeda et al. 2007). Artificial muscle-like actuators are receiving great interest for use as novel devices, because they are ideal for the realization of biomimetic movements by changing their shapes and dimensions. Since large deformation can be generated in CLCPs with the help of photochemical reaction of chromophores, light-driven soft CLCPs thus play an important role in realizing soft actuators, which can convert light energy into mechanical work directly by using a contactless laser beam.

Macroscaled Actuators Ikeda et al. reported the first light-driven plastic motor with laminated films composed of a CLCP film and a polyethylene sheet (Yamada et al. 2008). A continuous plastic belt of the CLCP laminated film was prepared by connecting both ends of the film and the belt placed on a homemade pulley system as illustrated in Fig. 15a. By irradiating the belt with UV light from top right and visible light from top left simultaneously, a rotation of the belt was induced to drive the two pulleys in a

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Fig. 15 (a) Schematic illustration of a light-driven plastic motor system used, showing the relationship between light irradiation positions and a rotation direction. (b) Photographs showing time profiles of the rotation of the light-driven plastic motor with the CLCP-laminated film induced by simultaneous irradiation with UV and visible light at room temperature (Yamada et al. 2008). (c) Series of photographs showing time profiles of the photoinduced inchworm walk of the CLCP laminated film by alternative irradiation with UV (366 nm, 240 mW cm2) and visible light (>540 nm, 120 mW cm2) at room temperature. The film moved on the plate with 1 cm  1 cm grid. Size of the film, 11 mm  5 mm; the CLCP laminated part, 6 mm  4 mm. Thickness of the layers of the film: PE, 50 μm; CLCP, 18 μm (Yamada et al. 2009)

counterclockwise direction at room temperature, as shown in Fig. 15b. Furthermore, they demonstrated a unidirectional motion, an inchworm walk, of the CLCP laminated film with asymmetric end shapes (Yamada et al. 2009). The film moved forward upon alternate irradiation with UV and visible light (Fig. 15c). Additionally, they also showed such creative 3D movements as a flexible robotic arm motion assembled with the CLCP laminated film. van Oosten et al. presented a new approach for producing cilia-like microactuators through inkjet printing followed by photopolymerization process (van Oosten et al. 2009). They used two dyes varying the composition of the actuator in the plane. The deposition with an inkjet printer allows different LC materials to be arranged perpendicular to the substrate at bottom and parallel to the substrate at the top of the film. The splayed molecular alignment in the film made the direction of the response independent from the direction of the incident light and generated stronger bending. The position of the actuator therefore could be brought into four positions by selecting the color composition of the light: When irradiated with only UV light, the yellow part of the actuator bent. When irradiated with both UV and visible light, the flap bent over its total length. Finally, if only visible light was used, only the red part of the flap bent. Switching between these four positions produced a cilia-like

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Fig. 16 (a) Artificial, light-driven cilia produce an asymmetric motion controlled by the spectral composition of the light. (b) Schematic representation of the macroscopic setup, showing the orientation of the molecules. (c) Steady-state responses of a 10-μm-thick, 3-mm-wide, and 10-mm-long modular liquid crystal network actuator to different colors of light (scale bar 5 mm). (d) Side view of the actuation of polymer cilia with ultraviolet light (1 W cm 2) in water (van Oosten et al. 2009)

motion (Fig. 16). This allowed a well-controlled motion by simply changing the color of the light. Therefore, this design overcame the inherent problems related to the use of electrical fields in wet environments. Moreover, the inject printing process allows fabrication of large-area and roll-to-roll active all-polymer devices and opens up possibilities for rapid prototyping of low-cost micromechanical systems. Instead of covalently linking the azobenzene moiety to the elastomer, PalffyMuhoray and coworkers created CLCP with an azo dye dispersed in it (CamachoLopez et al. 2004). When floating on water, the CLCP was found to swim into the darker regions, namely, away from the laser beam as a result of exchanging momentum between water and the sample upon its bending motion. Yu et al. prepared a visible light-driven fully plastic microrobot (Cheng et al. 2010a). The microrobot was made of crosslinked liquid crystalline polymer and PE bilayer films and consisted of several parts, including a hand, a wrist, and an arm (Fig. 17). Without the aid of any gears, bearings, or contact-based driving systems, the microrobot was manipulated to pick, lift, move, and place milligram-scale objects by irradiating different parts of the microrobot with visible light.

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Fig. 17 Schematic illustration of the microrobot and photographs showing microrobot picking, lifting, moving, and placing the object to a nearby container by turning on and off the light (470 nm, 30 mW cm2). The thickness of both PE and films was 20 μm. They were connected with each other by an adhesive. White arrows denote the parts irradiated with visible light (Cheng et al. 2010b)

Zhu et al. designed a model of photo-activated micropump, mainly including photodeformable material, pump membrane, pump chamber, and pipes (Chen et al. 2010). Water is chosen as the pump medium. The flow rate of the water varied in a stroke of the pump membrane, which means, the bending speed of the laminated film decreased in this process. The smaller pressure would lead to a higher flow rate and a larger volume pumped in a stroke. In further study, they utilized the bending of CLCP films to act as a valve membrane (Chen et al. 2011).

Microscaled Actuators Yu et al. designed and fabricated a light-regulated adhesion switch on a microarrayed azobenzene CLCP superhydrophobic surface, by which water droplets could be rapidly, precisely, locally, and through no contact controlled (Li et al. 2012b). It is well known that surface chemical composition as well as a suitable micro /nanoscale rough surface cooperatively creates superhydrophobicity. Polydimethylsiloxane-soft-template-based secondary replication was utilized to introduce uniform and quantitatively controllable surface roughness to the azobenzene CLCP film. After UV light irradiation, the azobenzene mesogens at

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Fig. 18 (a) Schematic illustration of the PDMS-soft-template-based secondary replication process. (b) Optical photo of microarrayed CLCP film with two patterned areas named as D15 and D5. Large-area optical microscopic image and local amplified image (inset). The patterns of D15 and D5 are all square-arrayed square posts with the post width of 10 mm. The spacings between two nearest posts for D15 and D5 are 15 mm and 5 mm, respectively. (c) Light-controlled quick and reversible switching of superhydrophobic adhesion between rolling and pinning on microarray CLCP with a 2 μL water droplet (Li et al. 2012b)

the surface of the film transferred to a cis state, and the water adhesion increased because of the growth of surface polarity, since in the excited cis-configuration, the dipole moment leads to an increase of local polarity of polymer chain. Subsequently, after the visible light irradiation, the azobenzene mesogens at the surface returned to trans state and the surface reverted to lower adhesive superhydrophobic state (Fig. 18). Such a quick and reversible switch of superhydrophobic adhesion was retained well after many cycles by the alternate irradiation of UV and visible light. Unlike contact angle (CA) switchable surface systems, our work puts emphasis on the switching of sliding angle (SA) on the same surface, while the static CAs before and after switching were all in the superhydrophobic range. Therefore, the “rolling” and “pinning” of water droplets was achieved, giving rise to promising applications in microfluids. It is the first time that the photoresponsive CLCP materials were used to prepare superhydrophobic adhesion switchable surfaces, which are of great importance for no-loss microdroplet transfer. Most recently, Yu et al. successfully fabricated CLCP films with different surface topographies and submicropillar and submicrocone arrays, through colloidal lithography technique by modulating different types of etching masks (Fig. 19, Zhan et al. 2015). The prepared submicropillar arrays were uniform with an average pillar diameter of 250 nm, and the cone bottom diameter of the submicrocone arrays was about 400 nm, which are much smaller than previously reported CLCP micropillars. More interestingly,

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Fig. 19 (a) Schematic procedure of the fabrication of microarrayed CLCP films. (b) SEM image of submicropillar arrayed CLCP film. (b) The shape of a water droplet on the CLCP film when it is turned upside down, indicating its high water adhesion. (c) SEM image of the surface of submicrocone arrayed CLCP film (Zhan et al. 2015)

these two species of films with the same chemical structure represented completely different wetting behavior of water adhesion and mimicked rose petal and lotus leaf, respectively. Both the submicropillar arrayed film and the submicrocone arrayed film exhibited superhydrophobicity with a water contact angle (CA) value of 144.0  1.7 and 156.4  1.2 , respectively. Meanwhile, the former demonstrated a very high sliding angle (SA) greater than 90 , and thus, the water droplet was pinned on the surface as rose petal. On the contrary, the SA of the submicrocone arrayed CLCP film consisting of micro- and nanostructure was only 3.1  2.0 , which is as low as that of lotus leaf. Compared to replica molding technique and inkjet printing technology used to fabricate microstructured CLCPs, colloidal lithography technique is timesaving and can be modulated throughout etching procedure, which finely regulates the structural parameters such as shapes and dimensions. Our work provides a new way to fabricate the CLCPs in the size of nanoscale.

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Yu et al. fabricated a photoresponsive 2D microarray with a period of about 1 μm using CLCP containing azobenzene groups by using the replica molding technique (Fig. 20, Yan et al. 2012). The CLCP microarray showed switchable behavior on the reflection spectra by alternate irradiation of UV and visible light, accompanying with the deformation of the CLCP pillars. Because the trans-cis photoisomerization results in the alignment change of LC molecules especially in the side surface region of pillars, the pillars contract along the long axis while they undergo expansion along the short axis, leading to the increase in the diameter of the pillars and variation in reflection spectra of the CLCP microarray. This is the first time to use the CLCP to fabricate the microarray with a period of about 1 μm and manipulate switchable behavior on the reflection spectra of the LC polymer microarray by light. Furthermore, Yu et al. prepared novel photo and thermal dual-responsive inverse opal films based on CLCP (Fig. 21, Zhao et al. 2014). The inverse opal film showed switchable behavior on the reflection spectra by alternate irradiation of UV and visible light or temperature, owing to the change in the order of the holes. This change in the periodic structure is ascribed to the contraction of CLCP induced by the photochemical reactions of the azobenzene moieties or the thermal-induced phase transition. The optical properties drastically decreased by thermal or photoinduced phase transitions of the CLCP. It is the first time that azobenzene-containing CLCPs have been used to prepare inverse opal film and achieve repeatable switching behavior on the reflection spectra of film by using light, which can be manipulated conveniently and controlled in situ. The reflectivity changed to a greater extent compared to the 2D CLCP photonic crystals. Wiersma et al. used direct laser writing system to pattern the complex 3D structures with sub-micrometer resolution (Zeng et al. 2014, 2015a). Microrobots were fabricated with the CLCPs acting as the main body of walkers (Zeng et al. 2015b). The light-induced maximum stress of these systems was measured to be 260  2 kPa, which was comparable to natural muscles (10–200 kPa). The legs of walkers had a conical shape, which was chosen to reduce the surface contact area, while 45 tilt of the leg created asymmetry adhesion necessary for walking (Fig. 22a and b). The artificial creature automatically performed various locomotion highly depended on the interactions with the environment. The microscopic walker finished random or directional walking, rotating, or jumping when placed on surfaces with different treating methods (Fig. 22c).

Summary It is the elegant combination of self-organization of properties of LCs and mechanical performance of organic polymers that enables photoresponsive CLCPs to directly convert light energy into mechanical work. Light is an ideal stimulus, for it can be localized, selective, and nondamaging and allows for remote activation and delivery of energy to a system. Due to their photocontrollable properties without any aid of other motors, gears, and wires, it is very convenient and attractive to reduce the size of the photo-driven CLCP actuators for their potential

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Fig. 20 (a) SEM images of the pillars of the azobenzene CLCP microarray before (left) and after (right) UV irradiation. (b) Schematic illustration showing the change in the geometry of the pillars of the azobenzene CLCP microarray. (c) Reflection spectra of the azobenzene CLCP microarray under the UV light irradiation (365 nm, 20 mW cm2, 15 min) and the following visible light irradiation (530 nm, 20 mW cm2, 5 min) with the angle of incidence of 60 (Yan et al. 2012)

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Fig. 21 SEM surface images of (a) the SiO2 opal film and (b) the CLCP inverse opal film and SEM cross-sectional images of (c) the SiO2 opal film and (d) the CLCP inverse opal film. The inset is the locally amplified image. Thickness of the inverse opal film is about 17 mm. SEM images of the

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Fig. 22 (a) SEM image of a microwalker lying upside down. Scale bar: 10 μm. (b) Side view of the microwalker with 500 nm leg tip shown in the inset. Scale bar: 10 μm. (c) Top row shows the initial state of microwalkers on different surfaces. Bottom row shows the microwalker randomly walking on the polyimide-coated glass surface, rotating with one leg stuck onto the polyimide-coated surface, walking with self-reorientation on the clean glass surface, walking in the direction determined by the grating groove pattern (vertical). Insets of the top row show the schematics of the surface (Zeng et al. 2015a)

ä Fig. 21 (continued) inverse opal film (e) before and (f) after UV light irradiation. The red regular hexagon and straight lines represent the arrangement of the holes before UV light irradiation; the green hexagon and lines represent the arrangement of the holes after irradiation with UV light. After UV light irradiation, the shape of the hexagon becomes irregular, and the straight lines have become curves. (g) Reflection spectra of the azobenzene CLCP inverse opal film under UV light irradiation (365 nm, 50 mW cm2, 5 min) and subsequent visible light irradiation (530 nm, 20 mW cm2, 15 min). (h) Reflection spectra of the inverse opal film as a function of temperature (Zhao et al. 2014)

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application in micro- and even nanoscales. So far, the microarrayed CLCP films, the inverse opal CLCP films, the cilia-like microactuators, and the microwalkers have been developed based on the replica molding, inject printing technique, and laser writing systems. In addition, the size of the photo-driven actuators has been successfully reduced to nanoscale by colloidal lithography technique. These make it advantageous to use photo-driven CLCP actuators in a wide range of potential application fields such as microfluid systems, micro-optomechanical systems (MOMS), and optical devices. However, further efforts are still needed to make them of value in real-life applications. For instance, the improvement of the energy conversion efficiency, fatigue resistance, and strength and the continuous production of CLCPs remain challenges for researchers. In addition, there is an urgent need to integrate the CLCPs into functional and sophisticated devices with other materials because of the difficulty of the photoresponsive CLCPs to serve as the whole smart devices in real applications.

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Photoactive Liquid Crystalline Polymers

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Contents Definition . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . Introduction . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . Representative Architectures of Photoactive LCP . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . Photoactive Liquid Crystalline Elastomers . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . Hyperbranched Architectures . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . Holographic Recording . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . Specific Cases . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . Holograms . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . Multiple Holograms in Single Holographic Medium . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . Chiral Photoactive Liquid Crystalline Polymer . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . Mechanical Motion in LC/Azo Polymers . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . Liquid Single Crystal Elastomers (LSCEs) . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . Opto-mechanical Effect in Liquid Single Crystal Elastomers . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . LC-Polymer Composite . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . Polymer-Stabilized Liquid Crystals (PSLC) . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . Polymer-Dispersed Liquid Crystals (PDLCs) . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . Miscellaneous . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . Photoactive Liquid Crystalline Anisotropic Gel . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . Photoactive Non-covalent LCP . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . Nanoparticle-Liquid Crystalline Elastomer Composites . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . Photoresponsive Main-Chain Oligomers . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . Macrocycles Containing Azobenzene Main-Chain Oligomers . . . . . . . . . . . . . . . . . . . . . . . . . . . . . Stimulus-Induced Deformations in Various Geometries . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . Ferroelectric Liquid Crystals . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . Light-Sensitive Microcapsules Based on Liquid Crystalline Polyester . . . . . . . . . . . . . . . . . . . . Photoresponsive Membranes . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . Future Directions . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . References . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . .

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A. B. Samui (*) Institute of Chemical Technology, Mumbai, India e-mail: [email protected] © Springer Nature Switzerland AG 2020 L. Zhu, C. Y. Li (eds.), Liquid Crystalline Polymers, Polymers and Polymeric Composites: A Reference Series, https://doi.org/10.1007/978-3-030-43350-5_56

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Abstract

Photopolymers make a distinct class of materials that have the ability to undergo physical and chemical change by the action of light. This has another important aspect in that the light stimulus can be applied from remote location without coming in contact with the substrate. At the outset of this chapter, introduction has been made on various combinations and architectures of the photoactive liquid crystalline polymers. The importance of photoresponsive LC polymers in holograms making and their important types are discussed. This is followed by discussion on photomechanical properties of photoresponsive LC polymers, composites, and elastomers. Individual mechanical change in multiple addressable segments of a system is discussed as this will enable future smart complicated functions. Liquid single crystalline elastomers with superior optomechanical properties have been discussed. Photoresponsive LC polymer composites comprising polymer-stabilized liquid crystals (PSLC) and polymerdispersed liquid crystals (PDLCs) have been briefly presented. This is followed by deliberations on the photoactive LC anisotropic gel and photoactive noncovalent LCP, such as liquid crystalline ionomers containing azobenzene mesogens and hydrogen-bonded liquid crystalline azo polymer, respectively. Nanoparticle-LC elastomer composites that have light-responsive nanoparticles and photoresponsive main-chain oligomers useful for studying various mainchain characteristics are elaborated. Then, the macrocycles containing azobenzene main-chain oligomers, useful in characterizing the structure–property relationship, and stimulus-induced deformations in various geometries of photoresponsive LCPs have been discussed. Ferroelectric liquid crystals, useful for electronic and photonic devices, are also added to the discussion. Light-sensitive microcapsules based on LCP and photoresponsive membranes useful in the area of biotechnology, including drug delivery, biosensing, microfluidics, lightpowered molecular machines, molecular shuttles, data storage, etc., have been elaborated. The chapter ends with future directions. Keywords

Photoactivity · Azo dye · Photoactive liquid crystalline polymer · Orientation · Holographic recording · Ionomer · Hydrogen bonding

Definition Photoactivity

Photoactive liquid crystalline polymers

There are several molecules in chemistry and biological system which on exposure to light of specific wavelength undergo physical or chemical changes. The liquid crystalline polymers (LCP) showing activity under radiation having specific wavelength are called photoactive liquid crystalline polymer.

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Example of photoactive moiety

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The photoactivity is attained either by introducing photoactive group in liquid crystalline polymer or the polymer inherently possess both the properties. The physical change such as isomerization of specific moiety is reversible. The chemical change such as crosslinking is usually irreversible or in some cases reversible. Azo, spiropyrans, -C=C-, etc.

Introduction Photo-controlled and photo-switchable materials based on liquid crystals (LC) are finding application in many areas such as colored information recording, photooptical triggers, and display units and more (Barth et al. 2005). It is well-known that optical and electrical properties are strongly associated with orientation and morphology of these polymers (Rajaram et al. 1996). Liquid crystal has shown increasing interest because of their self-assembling behavior and also natural ability to dramatically alter the properties of light reflected from or transmitted through the assembly. After the photoinduction phenomenon came to the knowledge of the researcher, it was quickly understood that incident light itself can control the molecular ordering or disordering of liquid crystalline system. This has become the basis for controlling system with light or “light control light” and the gate opened for high speed information processing. The photoactivity normally originates from photoactive groups having the ability to isomerize or in some cases form cross-links. Few common photochromic moieties with their photoactivity modes are described in Fig. 1. During UV radiation the original stable linear trans-isomer converts to bent Cisisomer, which disturbs the LC arrangement and phase transition occurs to iso-phase. The trans-isomer reappears after discontinuing the radiation, which is due to back isomerization triggering reassembly to LC phase (Fig. 2). On heating the LC structure crosses the LC-Isotropic (LC-iso) transition temperature. AS described above the phenomenon is similar during UV radiation. Either radiation or heating discontinued the original LC phase reappears. During cooling, the temperature goes below LC-iso transition temperature and LC phase reappears. Light is considered as most interesting and convenient among various stimulating sources, as it is a clean energy and can be manipulated conveniently and controlled precisely. This adds additional dimension to the system of phase transition trigger. The polymer system having both liquid crystalline moiety (rod-shaped; mesogen) and photoactive moiety gives numerous possibilities, e.g., photochromic group, such as azo can be chemically attached creating various architectures, or it can be dissolved in LC polymer (LCP) to make an azo doped structures. It can also be LC molecule is controlled by azo polymer. The other molecular designs can be based on

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H

O O

O

H

O

Ph

Red

Yellow OR

O

O

H

hν' or Δ

H Ph

O

hν

OR

O

O

OR

UV (>260 nm)

+

UV (85 wt%) in LCRBCs results into amorphous (crystallization temperature, Tc ~ 30 to 50  C) randomly embedded sphere-like PEO domains. In current work, we prepare LCBBCs containing electrical and magnetic responsive cyanobiphenyl derivative as the LC block and PEG as the brushlike block unit of composition wherein the PEG forms cylindrical domains within LC matrix. We compare topology and structure of these materials to the structure formed by its known linear counterparts. These materials will be useful for various stimuliresponsive applications including electrochemical devices and batteries.

Experimental Section Synthesis The monomers were synthesized as previously reported (Deshmukh et al. 2013a, 2014). A representative ring-opening metathesis polymerization (ROMP) of LCBBC73 is described as below (Fig. 1). In separate vials NBCB12 (0.8 g, 1.6 mmol) and NBMPEG (0.299 g, 0.14 mmol) with mG2nd (0.0315 mg, 0.043 mmol) are weighed. The vials are purged with nitrogen for 10 min after which anhydrous CH2Cl2 is added to each vial and stock solutions are prepared. Then, catalyst solution in CH2Cl2 is transferred to stirring NBCB12 monomer solution under inert atmosphere. After 15 min of polymerization of NBCB12, NBMPEG monomer solution is added under inert atmosphere. Polymerization of second block is allowed for another 15 min at room temperature and terminated by adding excess ethyl vinyl ether (EVE). The resulting polymer is precipitated in methanol and then filtered, followed by drying overnight under vacuum at room temperature.

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Thermal Analysis TA-2920 DSC (Q-100 series) instrument is used to analyze thermal properties of the copolymer. Indium standard is used to calibrate instrument. Roughly 5–10 mg of polymer sample is taken, and scanning rate of 10  C/min is used. Phase transition temperatures are determined by the first cooling or the second heating cycle using TA universal analysis software. Thermal gravimetric analysis (TGA) was performed on PerkinElmer TGA Q-500. The sample (6–10 mg) was heated from room temperature to 600  C at a ramping rate of 20  C/min under nitrogen flow. The TA universal analysis software was used to obtain the degradation temperature of the polymer.

Small-Angle X-Ray Scattering (SAXS) (Performed at University of Connecticut) SAXS is performed on a pin-hole collimated Rigaku SMAX3000 instrument configured with CuKα radiation (1.542 Å) produced by a micro-focus source. The beam diameter on the sample plane is 1 mm and the scattered intensity is recorded on a gas-wire electronic area (2D) detector. The area detector has a resolution of 1024  1024 pixels and is located at a distance of ~ 150 cm from the sample. The instrument is calibrated using Silver behenate standard with a d-spacing of 58.38 Å. SAXS data acquired is plotted as intensity (I) versus scattering vector q, where q = (4π/λ)  sin θ, (2θ is the scattering angle).

Results and Discussions Synthesis Two monomers, norbornenyl end functionalized PEG, where Mn of PEG chains = 2000 g/mol (NBMPEG) (Deshmukh et al. 2013) and cyanobiphenyl liquid crystalline units attached with 12 methylene spacer to norbornene (NBCB12) (Deshmukh et al. 2014), are synthesized according to previously reported procedures. Representative LCBBCs are synthesized by sequential ring-opening metathesis polymerization (ROMP) using a modified Grubbs catalyst second generation (mG2nd) ((H2IMes)-(pyr)2(Cl)2RuCHPh) (Love et al. 2002) using monomers NBMPEG and NBCB12 (Fig. 2). The compositions of LCBBCs are determined by integrating characteristic peaks of NBCB12 (a proton at 6.9 ppm) and NBMPEG (three protons at 3.36 ppm) in 1H NMR spectra and by comparing the ratio of their integration values (Fig. 2). The number average of molecular weight and polydispersity index of LCBBCs was determined using gel permeation chromatography (GPC) using THF as the eluent at room temperature. We will use the following

Fig. 2

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nomenclature: LCBBCx-25 k represents LC brush block copolymer, where x represents weight percentage (wt %) of NBCB12 and 25 k represents total MW of copolymers. The composition, molecular weight, and PDI of LCBBC are compiled in Table 1.

Thermal Properties The thermal gravimetric analysis (TGA) was used to study the thermal stability of the copolymer sample. The LCBBC73–25 k undergoes thermal degradation beginning (onset) at 370  C with a total mass loss of 4.9%. 50% of the copolymer degrades at 416  C. There is a small amount (8%) of inert residue remaining at 600  C (Fig. 3). Thermal properties of LCBBCs are investigated using differential scanning calorimetry (DSC) as shown in Fig. 4. Transition temperature and enthalpy values corresponding to (LCBBC73–25 k) are shown in Table 2. The copolymer (LCBBC73–25 k) is initially heated to 150  C, cooled to – 40  C, and reheated to 150  C. The first heating cycle is used to eliminate the thermal history of the sample. Distinct Tg and TCl transitions can be seen in second cooling cycle of DSC thermograms (Fig. 4) and presented in Table 2. Tg is a characteristic of norbornene backbone, while TLC represents liquid crystalline order.

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Table 1 Molecular characterization of LCBBC73–25 k Polymer LCBBC73–25 k

Mol. Wt. (Kg/mol)a Theor. GPC (PDI) 26 32 (1.10)

Mol (%) of LC 27

Mol %b NBCB12 74.3 (73)

NBMPEG 25.7 (27)

a

Determined by GPC with ELSD detector, where THF was used as eluent and polystyrene standards were used to construct a conventional calibration b Mol % of each monomer is calculated from the NMR spectrum of the polymer, where peaks at 3.330 and 6.930 are used to quantify composition of NBMPEG and NBCB12, respectively. Values in parentheses indicate concentration of monomer in feed

100

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Temperature (°C) Fig. 3 TGA results obtained for LCBBC73-25 k showing thermal degradation beginning at 370  C and with a total residual mass of 95.1%

Microstructural Analysis of LCBBC73–25 k Figure 5 represents small-angle X-ray scattering (SAXS) diffractograms of LCBBCs where two types of scattering reflections are noted, including (i) microphase segregation (q and its higher-order reflections) and (2) LC order (qLC). Microphase segregation: SAXS curve suggests microphase segregation (q) with higher-order peaks and is indicated by blue arrows as shown in Fig. 6. Primary peak and its higher-order reflections appearing at ratio q, √3 q, and √4 q implies that the system forms hexagonally packed PEG cylinders in the LC major matrix of cyanobiphenyl domains. The d-spacing is dBCP = 15.31 nm, while inter-cylinder distance is dCYL = 17.7 nm (dCYL = (4/3)1/2dBCP). Fig. 5 also shows LC scattering peak at qLC = 1.17 nm 1 corresponding to a smectic layer period of dLC = 5.42 nm. According to the thermal analysis, the system

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Fig. 4 First cooling DSC curve of LCBBC73-25 k with heating rate of 10  C/min exhibiting Tg at 11.5  C and Tcl at 74.4  C

Table 2 Thermal characterization of LCBBC73–25 k Tg ( C)a 11.5

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Fig. 6 Schematics of LCBBC. (a) LCBBC architecture where cyanobiphenyl mesogen is attached with backbone by 12 carbon methylene spacer and mPEG side chain with MW of 2000 g/mol. (b) Illustration of LCBBC architecture. (c) Proposed hierarchical structure of LCBBC73-25 k consisting of (i) hexagonally packed PLA cylinders inside the major LC matrix (ii) smectic LC mesogens with homogeneous anchoring to IMDS is expected (Deshmukh et al. 2014). 1D SAXS data supports the proposed hierarchical structure

features a stable liquid crystalline phase as confirmed by qLC peak appeared in SAXS profile. Twelve carbon spacers attaching norbornene backbone with LC domains allow sufficient freedom for LC domain arrangement and avoid entanglements with the backbone. Formation of smectic LC phase in LCBBC73–25 k suggests that 12 methylene spacer is long enough to permit sufficient motional decoupling between backbone and cyanobiphenyl LC mesogens. In summary, LCBBC73–25 k polymer having brush architecture with overall MW of 25 k shows (i) hierarchical structure consisting of microphase segregation suggesting hexagonally packed cylinders of PEG enclosed in LC domains with d-spacing of 15.31 nm and inter-cylinder distance of 17.7 nm and (ii) smectic ordering of LC domains due to 12 carbon spacer on the length scale of 5.42 nm. While we didn’t explore 2D SAXS and WAXS on this particular library, previous work using the same LC monomer but PLA as the other brushlike unit showed that the LC orients parallel to the intermaterial dividing surface (Deshmukh et al. 2014). In this particular library, we expect the LC molecules to orient parallel to IMDS with amorphous PEG chains present in the cylindrical domains. A cartoon representation of this hierarchical structure is shown in Fig. 6.

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Importance of Topology In bottlebrush diblock copolymers, two different variables control polymer composition, the length of the polymer backbone and length of the side chains. Blocks having symmetric brush lengths can be visualized as elongated soft cylinders, and increasing the block backbone lengths affects only in further elongation of soft cylinder, while it does not have pronounced effect on the cross-sectional area between the blocks, which is very important parameter for molecular packing. Repulsion of high-density side chains on each repeating unit of bottlebrush block copolymers leads to stretching of the backbone to obtain higher sizes of lamellar domains even at asymmetric backbone lengths of each block. Rzayev in his work studied the effect of asymmetric topology on the similar length blocks, comprising of longer side chain lengths on one block and shorter side chain lengths on the other blocks, on the change in cross-sectional area to get curved cylindrical phases (Rzayev 2009). Current system LCBBC73–25 k, with LC side chains on one block and semicrystalline PEG side chains with 2000 gm/mole MW on other block, though cannot be visualized as highly asymmetric bottlebrush because block with shorter length LC side chains can be considered analogous to linear block. Because of the linear nature of one block in LCBBC73–25 k polymer molecule, change in composition is obtained by varying the block lengths (ΦLC = 0.73) unlike asymmetric bottlebrush where similar block lengths are used with varying side chain lengths to obtain the difference in volume fractions. However, in LCBBC73–25 k, large asymmetry between PEG side chains (Mn = 2000 g/mol) and cyanobiphenyl side chains can be the driving force to obtain curved interphase cylindrical domains unlike bottlebrushes. As we have already studied the magnetic function of cyanobiphenyl LC to get magnetic field-based self-assembly in hexagonally packed cylindrical system with minor block forming the cylinders and major phase consisting of LC ordering (Deshmukh et al. 2014; Gopinadhan et al. 2014), it would be insightful to study the effect of PEG brushlike cylinder forming minor block as compared to linear PEG minor block in the system aligned with magnetic field-based self-assembly. Comparison in the systems with brushlike minor block with the linear minor block counterpart may shed light on the effect of brush architecture of minor block on d-spacing, inter-cylinder spacing, and order-disorder temperatures of phase segregation. Brush systems tend to phase segregate with larger domain size due to backbone stretching as compared to their linear counterparts. Larger aligned PEG domains would be feasible to have more metal ion loading for anisotropic transport as in case for the lithium-ion batteries. While we have not synthesized the block copolymer system with linear PEG forming minor block with side chain cyanobiphenyl LC forming the major phase, block copolymers with similar architectures were studied by Majewski et al. to obtain anisotropic ion transport in highly aligned hexagonally packed cylinders of linear block PEG after magnetic field-based self-assembly (Majewski et al. 2010). Magnetic functionality due to cyanobiphenyl mesogen is important because it

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provides efficient means to obtain highly aligned vertical nanostructures in thin film geometries on large area films. Magnetic fields were used to produce vertically aligned Li + conducting cylindrical domains made of linear block PEG in a block copolymer film over millimeter length scales. The architecture used in this particular study can be seen in Fig. 7. As shown in Fig. 7, six-carbon spacer in major block helps in motional decoupling between backbone and LC molecules to obtain coexistence of smectic LC phase of cyanobiphenyl mesogen within the block copolymer superstructures. At fPEO: 0.23, hierarchical structure consisted of hexagonally packed PEG cylinders inside the major LC matrix, where cyanobiphenyl LC assembles in smectic LC phase which are homogeneously anchored to IMDS. Amorphous PEG confined in microphase-segregated nanodomains at room temperature are of particular importance because they can potentially be used for preparing PEG-based lithium-ion batteries and other solid-state electrochemical devices. PEG crystallization was significantly suppressed by the inclusion of higher volume fraction of LC block, which in turn conducive for the anisotropic Li + ion transport in the highly aligned PEG cylinders. We previously explored the effect of increasing volume fraction of LC block on the suppression of PEG crystallization (Zhou et al. 2011). Majewski and coworkers also showed that complexation with Li+ ions results in further suppression of kinetically limited PEG crystallinity and so inclusion of Li+ did not play any significant role in increasing crystallinity of PEG. Suppression of crystallinity was consistent with well-resolved glass transition temperature around 56  C with Li+ loading up to PEG/Li+ 16:1. Temperature-dependent SAXS revealed that LC transition temperature (68  C, smectic-isotropic transition) was concurrent with the order-disorder transition temperature (TODT). This convergence of LC transition temperature with TODT of block copolymer architecture was attributed to low molar mass of a polymer. Temperature-dependent anisotropic conductivity of Li+ ions was studied, and it was suggested that the role of LC disorder in the temperaturedependent conductivity needed better understanding. Increasing the molecular weight of the polymer can decouple LC clearing temperature from the ODT to give more insight in this regard. LCBBC73–25 k system presents brushlike architecture PEG side chains in one block and LC side chains in the other block. Presence of this PEG brush architecture stretches the backbone of the brush block, and this can have impact on crosssectional area and domain spacing. Based on 1D SAXS studies, domain spacing of 15.3 nm is obtained for the block copolymer superstructures. DSC data suggests that the suppression of PEG crystallization is obtained which is quintessential to obtain

O m

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Fig. 7 Commercially available block copolymer sample, where minor block containing linear PEO and major block with poly methyl methacrylate backbone attached to cyanobiphenyl LC mesogen by six carbon spacer. Molecular characterizations: 10.4 kg/mol (PDI = 1.15), fPEO: 0.23

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amorphous PEG nanodomains for effective transport. Larger domain spacing for the hexagonally packed cylindrical PEG domains in brush-type architecture with the suppressed PEG crystallization gives better room for the higher loading of Li + loading. Decoupling of TODT and LC temperature by the virtue of higher molecular weight of a polymer can help in better understanding of the role of LC clearing temperature on the temperature-dependent anisotropic conductivity. Topological change due brush architecture of PEG side chains containing block can provide us following advantages over the block copolymer system with linear PEG block: (i) Suppression of PEG crystallization temperature for higher molecular weight polymer sample to obtain amorphous nanodomains conducive for anisotropic ion transport. (ii) Larger PEG domains due to stretching of the backbone due to PEO side chain crowding at higher molecular weights can help in higher loading of ions for transport. (iii) Presence of smectic LC structure within the block copolymer superstructure with amorphous PEG domains for ion transport. (iv) Higher molecular polymer system can be obtained to decouple TODT and LC clearing temperature which in turn can help in getting better understanding of the effect of LC clearing temperature on the temperature-dependent conductivity. Apart from these advantages, presence of vinyl groups in the norbornene backbone after polymerization provides sites for further cross-linking to obtain mechanically robust system with entrapped amorphous domains for the ion transport to realize the goal of solid-state electrolyte membranes. LCBBC73–25 system thus can be efficiently utilized to study the temperature dependence of anisotropic conductivity. It can help us to define parameter space to obtain magnetic field-based assembly of such brush-type polymers. Effect of field strength on the anisotropic conductivity at particular temperature can help us understand the practical relevance of these materials. Acknowledgments The authors thank Dr. Chinedum Osuji and Dr. Manesh Gopinadhan for their insightful comments. This work was supported by NSF under CMMI-1246804 and NSF under DMR-1507045.

References Bae J, Kim J-K, Oh N-K, Lee M (2005) Organization of Rigid Wedge-Flexible Coil Block Copolymers into liquid crystalline assembly. Macromolecules 38:4226–4230 Choo Y, Mahajan LH, Gopinadhan M, Ndaya D, Deshmukh P, Kasi RM, Osuji CO (2015) Phase behaviour of Polylactide-based liquid crystalline brush-like block copolymers. Macromolecules 48:8315–8322

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Deshmukh P, S-k A, Gopinadhan M, Osuji CO, Kasi RM (2013a) Hierarchically self-assembled photonic materials from liquid crystalline random brush copolymers. Macromolecules 46:4558–4566 Deshmukh P, Ahn SK, Ludovic GM, Kasi RM (2013b) Interplay between liquid crystalline order and microphase segregation on the self-assembly of side-chain liquid crystalline brush block copolymers. Macromolecules 46:8245–8252 Deshmukh P, Gopinadhan M, Choo Y, Ahn SK, Majewski PW, Yoon YS, Bakajin O, Elimelech M, Osuji CO, Kasi RM (2014) Molecular Design of Liquid Crystalline Brush-like Block Copolymers for magnetic field directed self-assembly: a platform for functional materials. Macro Lett 3:462–466 Gopinadhan M, Majewski PW, Osuji CO (2010) Facile alignment of amorphous poly(ethylene oxide) microdomains in a liquid crystalline block copolymer using magnetic fields:toward ordered electrolyte membranes. Macromolecules 43(7):3286–3293 Gopinadhan M, Deshmukh P, Choo Y, Majewski PW, Bakajin O, Elimelech M, Kasi RM, Osuji CO (2014) Thermally switchable aligned Nanopores by magnetic-field directed self-assembly of block copolymers. Adv Mater 26(30):5148–5154 Hamley IW, Castelletto V, Lu ZB, Imrie CT, Itoh T, Al-Hussein M (2004) Interplay between Smectic ordering and microphase separation in a series of side-group liquid-crystal block copolymers. Macromolecules 37:4798–4807 Hammond MR, Mezzenga R (2008) Supramolecular routes towards liquid crystalline side-chain polymers. Soft Matter 4:952–961 He X, Sun W, Yan D, Liang L (2008) Novel ABC2-type liquid-crystalline block copolymers with azobenzene moieties prepared by atom transfer radical polymerization. Eur Polym J 44:42–49 Jia L, Cao A, Levy D, Xu B, Albouy PA, Xing X, Bowick MJ, Li MH (2009) Smectic polymer vesicles. Soft Matter 5:3446–3451 Lee H, Pietrasik J, Sheiko S, Matyjaszewski K (2010) Stimuli-Responsive Molecular Brushes. Prog Polym Sci 35:24–44 Li M, Keller P (2006) Artificial muscles based on liquid crystal elastomers. Phil Trans R Soc A 364:2763–2777 Love JA, Morgan JP, Trnka TM, Grubbs RH (2002) A practical and highly active ruthenium-based catalyst that effects the cross metathesis of acrylonitrile. Angew Chem Int Ed 41:4035–4037 Mahajan LH, Ndaya D, Deshmukh P, Peng X, Gopinadhan M, Osuji CO, Kasi RM (2017) Optically active elastomers from liquid crystalline comb copolymers with dual physical and chemical cross-links. Macromolecules 50:5929–5939 Majewski PW, Gopinadhan M, Jang W, Lutkenhaus JL, Osuji CO (2010) Anisotropic ionic conductivity in block copolymer membranes by magnetic field alignment. J Am Chem Soc 132:17516–17522 Mao G, Ober CK (1997) Block copolymers containing liquid crystalline segments. Acta Polym 48:405–422 Rzayev J (2009) Synthesis of polystyrene-polylactide bottlebrush block copolymers and their melt self-assembly into large domain nanostructures. Macromolecules 42:2135–2141 Rzayev J (2012) Molecular bottlebrushes: new opportunities in nanomaterials fabrication. J ACS Macro Lett 1:1146–1149 Schneider A, Zanna J, Yamada M, Finkelmann H, Thomann R (2000) Competition between liquid crystalline phase symmetry and microphase morphology in a chiral smectic liquid crystallineisotropic block copolymer. Macromolecules 33(3):649–651 Tao YF, Zohar H, Olsen BD, Segalman RA (2007) Hierarchical nanostructure control in rod-coil block copolymers with magnetic fields. Nano Lett 7:2742–2746 Tenneti K, Chen XF, Li CY, Wan XH, Fan XH, Zhou QF, Rong LX, Hsiao BS (2008) Competition between liquid crystallinity and block copolymer self-assembly in core-shell rod-coil block copolymers. Soft Matter 4:458–461 Tenneti K, Chen X, Li CY, Shen Z, Wan X, Fan X, Zhou Q-F, Rong L, Hsiao BS (2009) Influence of LC content on the phase structures of side-chain liquid crystalline block copolymers with bent-Core Mesogens. Macromolecules 42:3510–3517

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Verploegen E, Zhang T, Jung YS, Ross C, Hammond PT (2008) Controlling the morphology of side chain liquid crystalline block copolymer thin films through variations in liquid crystalline content. Nano Lett 8:3434–3440 Wu B, Mu B, Wang S, Duan J, Fang J, Cheng R, Chen D (2013) Triphenylene-based side chain liquid crystalline block copolymers containing a peg block: controlled synthesis, microphase structures evolution and their interplay with discotic mesogenic orders. Macromolecules 46:2916–2929 Xu B, Piñol R, Nono-Djamen M, Pensec S, Keller P, Albouy P-A, Lévy D, Li M-H (2009) Selfassembly of liquid crystal block copolymer PEG-b-smectic polymer in pure state and in dilute aqueous solution. Faraday Discuss 143:235–250 Yang J, Piñol R, Gubellini F, Lévy D, Albouy P-A, Keller P, Li M-H (2006) Formation of polymer vesicles by liquid crystal amphiphilic block copolymers. Langmuir 22:7907–7911 Yu H, Kobayashi T (2009) Fabrication of stable Nanocylinder arrays in highly birefringent films of an amphiphilic liquid-crystalline Diblock copolymer. ACS Appl Mater Interfaces 1:2755–2762 Zhang M, Müller AH (2005) Cylindrical polymer brushes. J Polym Sci A Polym Chem 43:3461–3481 Zhou Y, Ahn S-K, Lakhman RK, Gopinadhan M, Osuji CO, Kasi RM (2011) Tailoring crystallization behavior of PEO-based liquid crystalline block copolymers through variation in liquid crystalline content. Macromolecules 44:3924–3934

Photoalignment of Liquid Crystal Molecules Using Fluorine-Containing Polyimides

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Shuichi Sato, Hironaga Matsumoto, Setsuko Matsumoto, and Kazukiyo Nagai

Contents Definition . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . Photoalignment of Liquid Crystal Molecules . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . Substituent Effect of Fluorine-Containing Polyimide on Photoalignment of LC Molecules . . . Photoreactivity of 6FDA-TeMPD Polyimide Prepared with Various Film Preparation Protocols as an LC Photoalignment Film . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . Photoalignment Characteristics of CT Complex-Type Polyimides with Tetramethyl Phenylene Diamine . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . Conclusions . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . References . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . .

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Liquid crystal (LC) alignment methods are very important for manufacturing liquid crystal displays (LCDs). The photoalignment method in which polyimide films are radiated with linearly polarized ultraviolet (LPUV) light is one of the most effective non-rubbing processes to solve problems such as electrostatic charge and dust accumulation. This entry presents a review on photoalignment

S. Sato Department of Electrical and Electronic Engineering, Tokyo Denki University, Adachi-ku, Tokyo, Japan H. Matsumoto Department of Electronics and Bioinformatics, Meiji University, Kawasaki, Kanagawa, Japan S. Matsumoto Department of Physics, Meiji University, Kawasaki, Kanagawa, Japan K. Nagai (*) Department of Applied Chemistry, Meiji University, Kawasaki, Kanagawa, Japan e-mail: [email protected]; [email protected] © Springer Nature Switzerland AG 2020 L. Zhu, C. Y. Li (eds.), Liquid Crystalline Polymers, Polymers and Polymeric Composites: A Reference Series, https://doi.org/10.1007/978-3-030-43350-5_68

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film of LC using new fluorine-containing polyimides. We have provided detailed spectroscopy study to understand the possible mechanism behind the photoalignment effect for liquid crystal molecules. Keywords

Liquid crystal · Photoalignment · Fluorine-containing polyimide · Charge transfer complexes · Optical property · Molecular orbital calculation

Definition LC alignment methods are very important for manufacturing LCDs. The alignment film located outside the LC layer is crucial for LCDs, because it is necessary to precisely control the movement of LC molecules. The photoalignment method in which polyimide films are radiated with LPUV light is one of the most effective non-rubbing processes. We have summarized new photoalignment polyimides that are different from the existing photoalignment polymers such as azo- and coumarincontaining polymers. We theorized that charge transfer interaction between the polymer segments strongly depended on the photoreaction caused by LPUV irradiation.

Photoalignment of Liquid Crystal Molecules LCDs are widely used in cell phones, computers, and televisions because of their high definition image combined with thinness and low power consumption compared with previously used cathode-ray tube displays. Figure 1 shows a cross section of an LCD. In general, LCDs consist of a light source, two pieces of a polarizing filter, an LC cell, a color filter controlling the color of the picture, and an antireflection film or the like. The LC cell is composed of an LC and LC alignment film located outside the LC layer so that the LC molecules can be oriented. The thickness of the LC alignment film is approximately 100 nm. The thin alignment film is crucial for LCDs, because it is necessary to precisely control the movement of LC molecules. In addition, a precise molecular design is required for the alignment film. The function and principle of an LCD are shown in Fig. 2. The LCD consists of two polarizing filters and LC molecules between two electrodes. The change in the optical refractive index when the voltage is applied controls the switching (ON/OFF). As is evident from Fig. 2, the LC alignment on the substrate is very important for manufacturing LCDs. It is generally impossible to orient LC molecules on the surface of a nontreated substrate. However, it is possible to orient them when a polymer film is coated on the surface of the substrate and some surface treatment is applied. Mechanical rubbing method has been extensively used in the alignment of LC molecules. However, mechanical rubbing causes problems such as electrostatic charge and dust accumulation (Cognard 1982). To solve these problems, the

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Fig. 1 Photograph and crosssection images of an LCD display

Fig. 2 LC director and polarizer film configurations of a normal LCD in the voltage-off state (left) and voltage-on state (right)

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development of non-rubbing processes is essential. Therefore, photoalignment of LC molecular films has received considerable attention. This is because photoalignment process does not depend on electrostatic charge and dust in rubbing process. The photoalignment method in which polyimide films are radiated with LPUV is one of the most effective non-rubbing processes (Schadt et al. 1992). Ha et al. and Shitomi et al. reported that the photoalignment characteristics of polyimide films, including alicyclic and aromatic compounds, clearly depend on the wavelength of the exposure photon (Shitomi et al. 1999; Ha and West 2002). Gibbons et al. showed that curing the orientation of photosensitive layers with polarized light may result in an “in-plane” orientation of the LCs on azo polymers (Gibbons et al. 1991). In addition, Schadat et al. observed LC photoalignment on coumarin-containing polymer layers (Schadt et al. 1992). However, these polymer materials have low heat resistance, low optical transparency, and are dark brown or yellowish brown. Our preliminary experiments showed that 4,4-(hexafluoroisopropylidene) diphthalic anhydride 2,3,5,6-tetramethyl-1,4-phenylene diamine (6FDA-TeMPD) exhibits photoalignment characteristics (Sato et al. 2011a). This fluorine-containing polyimide has high heat resistance (glass transition temperature, Tg = 427
C), high solubility in polar solvents, high optical transparency, and low refractive index compared with other polymers. This entry focuses on the photoalignment of LC molecules using fluorinecontaining polyimides.

Substituent Effect of Fluorine-Containing Polyimide on Photoalignment of LC Molecules 6FDA-containing polyimides are expected to show high heat resistances, low dielectric constants and refractive indices. The fluorine atom has high electronegativity, whereas the C-F bond has high binding energy and low polarizability. Thus, fluorine-containing polyimides have high potential as materials for use in the LCD industry. Most importantly, 6FDA is commercially available as a tetracarboxylic dianhydride component, leading to its widespread use. In general, polyimides containing 6FDA have excellent solubility in polar solvents and high optical transparency, along with a low refractive index. However, LC molecule photoalignment on 6FDA-based polyimides has not been studied. Hence, in the current study, we systematically investigated the substituent effect of fluorine-containing polyimides with a 6FDA group on the photoalignment of LC molecules and the effect of LPUV irradiation with a wavelength of 254 nm on the chemical structure and the optical properties in infrared (FTIR) spectra, ultraviolet (UV) absorption spectra, and molecular orbital (MO) calculations. In addition, the LPUV photoreaction was also investigated. The chemical structure of each product, as shown in Fig. 3, was confirmed by FTIR spectroscopy and nuclear magnetic resonance analyses. The 6FDA-based

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Fig. 3 Chemical structures and calculated models of fluorine-containing polyimides. 6FDA-based polyimides: 6FDA-mPD, 6FDA-MPD, 6FDA-TMPD, and 6FDA-TeMPD. (Reprinted with permission from J. Photopolym. Sci. Technol. 24, 617. (Copyright (2011) The Society of Photopolymer Science and Technology))

polyimides were 6FDA-1,3-phenylene diamine (mPD), 6FDA-4-methyl-1,3-phenylene diamine (MPD), 6FDA-2,4,6-trimethyl-1,3-phenylene diamine (TMPD), and 6FDA-TeMPD. The 6FDA-based polyimides used in the current research were the same samples that we used in a previous study (Miyata et al. 2008; Sato et al. 2010, 2011a, b, 2013). We fabricated isotropic, dense, and nonporous polyimide films using the spincoating method. Then, LPUV irradiation was performed with a 200 W Hg-Xe lamp light source. The measured polarization degree of the light at 254 nm was 99%. Irradiation intensity was about 700 μW/cm2. LC cells were fabricated by assembling two indium tin oxide substrates, with their irradiated LPUV light polarizations set parallel to each other according to the literatures (Nishikawa and West 1999; Shitomi et al. 1999). The cell gap was about 25 μm. The guest-host LC was mixed with ZLI-2293 LC and dichroic dye G-256, heated to 100
C, and then finally injected into the cells in an isotropic phase.

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The glass transition temperatures of the isotropic, dense, and nonporous 6FDAmPD, 6FDA-MPD, 6FDA-TMPD, and 6FDA-TeMPD polyimide films were 298, 335, 380, and 427
C, respectively (Miyata et al. 2008; Sato et al. 2010, 2011b). The glass transition temperature was much higher than the temperature used in all of the optical measurements. Under optical measurement conditions, all of the 6FDA-based polyimides were glassy and completely amorphous. As the number of methyl groups in the diamine moiety increased, the glass transition temperature also increased. The addition of four methyl groups led to an increase of 129
C. Thus, the methyl groups seemingly restricted the mobility of the polymer segments. The IR spectra of the LPUV-irradiated 6FDA-based polyimides for 2 J/cm2 and unirradiated 6FDA-based polyimides are shown in Fig. 4. 6FDA-mPD: 1786 cm1 and 1726 cm1 (C=O stretching), 1360 cm1 (C–N stretching), and 1346 cm1 (C–H bending). 6FDA-MPD: 1786 cm1 and 1729 cm1 (C=O stretching), 1360 cm1 (C–N stretching), and 1345 cm1 (C–H bending). 6FDA-TMPD: 1786 cm1 and 1728 cm1 (C=O stretching), 1358 cm1 (C–N stretching), and 1345 cm1 (C–H bending). 6FDA-TeMPD: 1786 cm1 and 1726 cm1 (C=O stretching), 1354 cm1 (C–N stretching), and 1342 cm1 (C–H bending). In Fig. 4a, the IR spectra of 6FDA-mPD, 6FDA-MPD, and 6FDA-TMPD did not change through LPUV irradiation. However, for 6FDA-TeMPD, the O-H bond peak intensity near 3500 cm1 slightly increased, the C=O bond peak intensity near 1730 cm1 decreased, and the peak intensity near 1360 cm1 decreased upon LPUV irradiation. The C-N and C–H bending peaks were observed in the range 1300–1400 cm1. The IR spectra in this range are shown in Fig. 4b. The peak whose intensity decreased upon LPUV irradiation is due to the C–H bending of the methyl group in the diamine moiety. The result indicates that only the 6FDA-TeMPD chemical structure changes with excitation of the C=O double bond and C-H bond. In addition, the hydroxyl group O-H bond was newly formed. However, no photoscission reaction of LPUV irradiation was observed, because the other peak intensities did not change. Therefore, in the current study, the LPUV reaction was systematically studied through its UV absorption and MO calculation. The UV spectra from 3.5 to 6.5 eV of these LPUV-irradiated and unirradiated 6FDA-based polyimides are shown in Fig. 5. Two significant peaks were observed in the ranges 5.5–5.7 eV (No. 1) and 6.2–6.4 eV (No. 2) in 6FDA-based polyimides. The No. 1 peak position shifted to lower energy as the number of methyl groups in the diamine moiety increased. The No. 1 peak positions were 5.68 eV for 6FDA-mPD, 5.62 eV for 6FDA-MPD, 5.53 eV for 6FDA-TMPD, and 5.52 eV for 6FDA-TeMPD. In contrast, the No. 2 peak positions did not change. The No. 2 peak positions were about 6.5 eV for 6FDA-mPD and 6FDA-MPD, 6.21 eV for 6FDA-TMPD, and 6.21 eV for 6FDA-TeMPD. The peak position remained constant and did not depend on the number of methyl groups in the diamine moiety. The photon energy associated with a wavelength of 254 nm is 4.9 eV. This photon energy was absorbed in all 6FDA-based polyimides used in the current study. A change in peak intensity was not observed in this photon energy range. However, the

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Fig. 4 FT-IR spectra of LPUV-irradiated 6FDA-based polyimide films from 0 to 2 J/ cm2. FT-IR spectra range: (a) 1000–4000 cm1, (b) 1300–1400 cm1. (Reprinted with permission from J. Photopolym. Sci. Technol. 24, 617. (Copyright (2011) The Society of Photopolymer Science and Technology))

Nos. 1 and 2 peak intensities of 6FDA-TeMPD clearly decreased upon LPUV irradiation. This tendency is in good agreement with that in the IR spectra shown in Fig. 4. This result indicates that the chemical structure of 6FDA-TeMPD was changed by LPUV irradiation. No photoscission reaction by LPUV irradiation was observed because the peak intensity at 4.9 eV did not decrease. Thus, these UV absorption spectra will be investigated later with the result of MO calculations. The calculated UV absorptions of 6FDA-based polyimides optimized in Fig. 3 are shown in Fig. 6. A peak position difference between the observed and calculated

500 Fig. 5 Measured absorption spectra of LPUV-irradiated 6FDA-based polyimide films from 0 to 2 J/cm2. (Reprinted with permission from J. Photopolym. Sci. Technol. 24, 617. (Copyright (2011) The Society of Photopolymer Science and Technology))

Fig. 6 Calculated absorption spectra of 6FDA-based polyimides. (Reprinted with permission from J. Photopolym. Sci. Technol. 24, 617. (Copyright (2011) The Society of Photopolymer Science and Technology))

S. Sato et al.

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absorption spectra was observed because the calculated absorption spectra were for one unit of their polymers. However, the existence of two peaks in the calculated absorption region was in good agreement with that in the observed absorption in the current study. Two significant peaks were observed from 5.6 to 5.8 eV (No. 1) and at 6.5 eV (No. 2) in the calculated 6FDA-based polyimides; the peak at 4.9 eV was absorbed. The No. 1 peak positions were 5.69 eV for 6FDA-mPD, 5.66 eV for 6FDA-MPD, 5.67 eV for 6FDA-TMPD, and 5.70 eV for 6FDA-TeMPD. These peak positions were almost constant. However, the calculated peak position was higher than the observed position. In contrast, the No. 2 peak position did not change. The calculated peak position was higher than the observed position, similar to the No. 1 peak. Therefore, in 6FDA-TeMPD, these two peaks (i.e., Nos. 1 and 2) changed by LPUV irradiation were assigned based on the result of the MO calculation. A comparison between the calculated and observed UV absorptions of 6FDATeMPD polyimide is shown in Fig. 7. The spectrum has two distinct peaks around 5.7 eV (No.1) and 6.5 eV (No. 2). Although the two peaks are broad and appear not to have fine electronic structure, the spectrum can be deconvoluted with various components considering the possible electronic transitions. The electronic absorption corresponds to the transition from the ground to the first excited state. It is mainly described by one electron excitation from the highest occupied MO (HOMO) to the lowest unoccupied MO (LUMO). Consequently, the HOMO ! LUMO transition implies electron density transfer. The 6FDA-TeMPD orbital compositions of the frontier MO are shown in Fig. 8. Note that the charges of the HOMO and Fig. 7 Measured and calculated absorption spectra of 6FDA-TeMPD. (Reprinted with permission from J. Photopolym. Sci. Technol. 24, 617. (Copyright (2011) The Society of Photopolymer Science and Technology))

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Fig. 8 Aromatic orbital HOMO and LUMO compositions of the frontier MO for 6FDA-TeMPD. (Reprinted with permission from J. Photopolym. Sci. Technol. 24, 617. (Copyright (2011) The Society of Photopolymer Science and Technology))

LUMO are localized on the TeMPD and 6FDA site residues, respectively. This finding indicates that the charge transfer (CT) can take place through the HOMO, HOMO-1 ! LUMO, LUMO+1 transitions, as well as through general polyimide. That is, photoabsorption depends on the CT transitions from the TeMPD site to the 6FDA site because of the phenyl group π–π transitions. The components can be provisionally assigned based on the electronic transition of a simple aromatic or imide compound. The band lower than 5.7 eV (No.1) is assignable to the π–π transition of the phenyl group and the bands at 5.7 eV (No.1) and 6.5 eV (No. 2) to the π–π transition of the phenyl group and n–π transition of the carbonyl groups (i.e., HOMO–5 ! LUMO or LUMO+1), respectively. As is evident in Fig. 5, the reduction of the only n–π transition of the carbonyl groups depended on the reduction of the photoabsorption of Nos. 1 and 2 peaks and not on those lower than 5.7 eV (No. 1). This result indicates that no photoscission reaction caused by LPUV irradiation occurred, which is in good agreement with the IR spectra. The formation of CT complexes (CTCs) or CT interaction is well-known in general polyimides. This interaction is attributed to the presence of π electrons between the ring structures in polyimides. One of the CTC structures shows mixed-layer packing (MLP), which is similar to a sandwich structure between imides and the neighboring benzene rings in the diamine moiety shown in Fig. 9a. The MLP can explain 6FDA-TeMPD polyimide films (Kanehashi et al. 2011). A large number

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Fig. 9 Mechanism for photooxidation of 6FDA-TeMPD. (Reprinted with permission from J. Photopolym. Sci. Technol. 24, 617. (Copyright (2011) The Society of Photopolymer Science and Technology))

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of aromatic polyimides showing various properties have been systematically synthesized using various chemical structures. The difference in polyimide molecular ordering seems to be closely related to their heat resistance property (i.e., Tg = 427
C). Phenyl group π–π transitions between polymer segments are expected to occur easily. Therefore, in Fig. 7, the observed Nos. 1 and 2 peak positions shifted to lower energy compared with the calculated position. The CT interaction between polymer segments strongly depended on the photoreaction upon LPUV irradiation. This result indicates that the chemical structure changes with the excitation of the C=O double bonds and C-H bonds and the formation of hydroxyl group O-H between the intermolecular polymer segments through the LPUV irradiation mechanism of 6FDA-TeMPD, as schematically shown in Fig. 9. The CTC structure of 6FDATeMPD was MLP, which is similar to the sandwich structure presented in a previous study (Kanehashi et al. 2011). The electron-donating benzene ring and the electronaccepting imide ring formed CTCs between both sides of the polymer chains in this MLP structure (Fig. 9a). Therefore, the photoreaction between the C=O group of the imide ring and a CH3 group upon LPUV irradiation occurred in one direction (Fig. 9b); hydroxyl group O-H was newly formed on the side chain. In the case of polymers with a symmetric structure, such as 6FDA-TeMPD, the combinations with alternating sequences to both cross sides clearly formed (Fig. 9c). This result is in a good agreement with that of the IR and UV spectra shown in Figs. 4 and 5, respectively. The polarization microscope (POM) images of all 6FDA-based polyimides in the current study are presented in Fig. 10. The photoalignment characteristic of all LPUV unirradiated 6FDA-based polyimides was not observed. The photoalignment characteristic of LPUV irradiation at 10 J/cm2 was also not observed in 6FDA-mPD, 6FDAMPD, and 6FDA-TMPD, as well as in the LPUV unirradiated 6FDA-based polyimide. However, in the case of 6FDA-TeMPD, the photoalignment characteristic of LPUVirradiated polyimides was observed. The photoalignment characteristic of LPUVirradiated 6FDA-TeMPD was estimated to depend strongly on the photoreaction between the C=O group of the imide ring and a CH3 group for the phenyl group π–π transitions based on Fig. 9. The CTCs easily formed in polymers with a large substituent, alternating sequence, and polarizability, such as 6FDA-TeMPD. Because of the photoreaction in the CTC structure, only 6FDA-TeMPD showed photoalignment of the LC molecule. However, 6FDA-mPD, 6FDA-MPD, and 6FDATMPD, which form CTC structures with difficulty, did not show photoalignment of the LC molecules. The chemical and physical structures of the polymer and polymer cohesiveness strongly affected the photoalignment of the LC molecules.

Photoreactivity of 6FDA-TeMPD Polyimide Prepared with Various Film Preparation Protocols as an LC Photoalignment Film Several studies have investigated the effects of different preparation conditions on the physical chemistry of condensation polymerization. However, the effects of the preparation protocol on the photoreactivities and LC alignments of

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Fig. 10 POM images of LPUV-irradiated 6FDA-based polyimide films from 0 to 10 J/cm2. (Reprinted with permission from J. Photopolym. Sci. Technol. 24, 617. (Copyright (2011) The Society of Photopolymer Science and Technology))

photoalignment films have not been reported to date; only the effects of the chemical structure on the LC alignment have been investigated. The film preparation protocols strongly depend on photoreactivity because of the photoreactivity dependence on both the residual casting solvent molecules and the condensation of the polymer segments in the films. However, no direct evidence for this hypothesis has been presented in the literature. Therefore, identifying the crucial factors that determine the LC alignment characteristics through various film preparation protocols is necessary in the development of high-performance photoalignment films. In the current study, 6FDA-TeMPD polyimide was used as a model polymer. To date, no systematic study on the dependence of photoalignment characteristics on film preparation conditions has been reported. Polymer chains exhibit a variety of conformations in different casting solvents. These different conformations result in different polymer film morphologies, including chain packing (Kanehashi et al. 2007, 2011). Thermal treatment induces the densification of glassy polymer films. Hence, the photoalignment mechanisms of 6FDA-TeMPD polyimide films prepared in different casting solvents were investigated. The effects of the different film preparation protocols on the heat-treated films were also analyzed. The 6FDA-TeMPD polyimides used in the current research were the same as the samples used in a previous study (Sato et al. 2011). The isotropic, dense polyimide

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films were prepared by spin-coating a filtered 2.0 wt% solution of the polymer in high volatility (boiling point = 67
C) tetrahydrofuran (THF) or low volatility (boiling point = 166
C) N,N-dimethylacetamide (DMAc) on quartz substrates at 3000 rpm for 30 s. The 6FDA-TeMPD polyimides were then baked at 100 or 250
C for 3 h. The resulting substrates were then allowed to cool to room temperature. The UV spectra (from 3.5 eV to 6.5 eV) of the films LPUV-irradiated at 10 J/cm2 as well as those of unirradiated 6FDA-TeMPD polyimide films are shown in Fig. 11. Three significant peaks, from 4.1 eV to 4.2 eV (peak a), 5.3 eV to 5.6 eV (peak b), and from 6.2 eV to 6.4 eV (peak c), were observed for all of the films. The intensity of peak a slightly increased after LPUV irradiation, whereas the intensities of peaks b and c clearly decreased. Moreover, peak c eventually disappeared. These results are consistent with those of a previous study (Sato et al. 2011). The position of peak a, which is the maximum absorption wavelength, λa max, for all films was found from 4.14 eV to 4.16 eV. However, that of peak b, λb max, depended on the film preparation protocol. In the high volatility THF solvent, the peak position slightly shifted toward the 5.50–5.54 eV range at 100
C and 5.51–5.55 eV range at 250
C during LPUV irradiation. In contrast, in the low volatility DMAc solvent, λb max shifted toward the 5.37–5.45 eV range at 100
C and 5.44–5.55 eV range at 250
C. The shift in λb max of the 6FDA-TeMPD film prepared in the DMAc solvent was more significant than that prepared in THF, indicating that the transition-state structures of the films were determined by the film preparation protocols.

Fig. 11 Absorption spectra of LPUV-irradiated 6FDATeMPD films from 0 to 10 J/ cm2. (Reprinted with permission from J. Photopolym. Sci. Technol. 25, 401. (Copyright (2012) The Society of Photopolymer Science and Technology))

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The frontier MOs of peaks a, b, and c were analyzed in our previous study (Sato et al. 2011). The 6FDA-TeMPD orbital compositions of the frontier MOs of these peaks are shown in Fig. 12. Peak a is mainly described by the excitation of one electron from the HOMO to the LUMO. The lowest band can be assigned to the π–π transition of the phenyl group. However, the bands of peaks b and c can be deconvoluted into various components by considering the possible electronic transitions. Both peaks a and b showed a dependence on the n–π transition of the carbonyl groups (HOMO-5 or HOMO-7 ! LUMO, respectively). Figure 11 shows that the reduction of only n–π transition of the carbonyl groups depended on the photoabsorption of peaks b and c. This reduction was in turn determined by the 6FDA-TeMPD film preparation protocols. In this study, peak c was not observed after LPUV irradiation. Therefore, only the intensities of peaks a and b, which decreased with increasing irradiation time, were analyzed. The peak intensity ratios of the films prepared via different protocols can be expressed as follows:

Fig. 12 Aromatic HOMO and LUMO compositions of the frontier MO of 6FDA-TeMPD. (Reprinted with permission from J. Photopolym. Sci. Technol. 25, 401. (Copyright (2012) The Society of Photopolymer Science and Technology))

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Peak intensity ratio ¼

Aa,yJ  Aa,0J Ab,yJ  Ab,0J , Aa,0J Ab,0J

(1)

where Aa, yJ and Ab, yJ are the absorbances of peaks a and b, respectively, under y J/cm2 LPUV irradiation, whereas Aa, 0J and Ab, 0J are the absorbances of peaks a and b, respectively, without LPUV irradiation. The peak intensity ratios of peaks a and b of the films prepared via different film preparation protocols as a function of the LPUV irradiation time are shown in Fig. 13a, b, respectively. The intensity of peak a for all 6FDA-TeMPD films Fig. 13 Peak intensity ratio for (a) peak a and (b) peak b in the absorption spectra as a function of LPUV exposure from 0 to 10 J/cm2. Film preparation conditions: THF 100
C (○), THF 250
C (●), DMAc 100
C (□), DMAc 250
C (■). (Reprinted with permission from J. Photopolym. Sci. Technol. 25, 401. (Copyright (2012) The Society of Photopolymer Science and Technology))

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increased by 0–6 J/cm2 as the irradiation time increased. However, the intensity decreased by 6–10 J/cm2 as the irradiation time further increased. The extents of the increase in the range 0–6 J/cm2 ranked as follows: THF 100
C > DMAc 250
C > THF 250
C > DMAc 100
C. In addition, the extents of reduction in the range 6–10 J/cm2 ranked as follows: THF 250
C > DMAc 250
C > DMAc 100
C > THF 100
C. The peak intensity was nearly constant because it did not depend on the temperature and casting solvent conditions. Meanwhile, the intensity of peak b decreased as the LPUV irradiation time increased. The reduction was in the order, THF 250
C > THF 100
C > DMAc 250
C > DMAc 100
C, and showed greater dependence on the casting solvent than on the heat-treatment temperature. This reduction also depended on the highly volatile preparation conditions. The peak intensity of the 6FDA-TeMPD films prepared in the DMAc solvent was not affected by irradiation at 1 J/cm2 LPUV (Fig. 11). In addition, a time lag in the reduction was observed. Therefore, the film preparation conditions affected the LPUV irradiation effect. The photoreaction mechanisms of all 6FDA-TeMPD films prepared under LPUV irradiation are schematically shown in the absorbance spectra (Fig. 9). The electron-donating benzene ring and the electron-accepting imide ring formed CTCs between the polymer chains in the MLP structure. Therefore, the photoreaction between the C=O group of the imide ring and the CH3 group under LPUV irradiation occurred in one direction, resulting in the formation of hydroxyl group O-H on the side 6FDA-TeMPD polymer chain. The symmetrically structured polymer (6FDA-TeMPD) exhibited combinations with alternating sequences on both cross sides. The increase in the intensity of peak a was determined by CTC formation, which is a measure of polymer cohesiveness. On the other hand, the reduction in the intensity of peak b is attributed to the decrease in the photoreaction of the C=O group. Therefore, the fluorescence spectra of all 6FDA-TeMPD films were obtained to characterize the CT interactions and determine the extent of CTC production, given that LPUV irradiation affects the condensation of polymer segments. Figure 14 shows the fluorescence spectra of all 6FDA-TeMPD films irradiated at 0, 1, and 10 J/cm2 LPUV. All 6FDA-TeMPD polyimide films exhibited an emission spectral band at 470 nm after excitation at 325 nm. This band is attributed to the intramolecular interactions of the aromatic polyimide that contains an alternating sequence of electron-rich donor and electron-deficient acceptor molecules. Therefore, CTC formation and polymer cohesiveness can be determined via fluorescence spectroscopy. The maximum peak emission wavelength, λ max, was 471 nm for the films thermally treated at 100
C (THF 100
C and DMAc 100
C), while λ max of the films thermally treated at 250
C (THF 250
C and DMAc 250
C) was 474 nm. The λ max depended on the thermal treatment temperature. In addition, the λ max of all prepared films increased to approximately 478 nm as the LPUV irradiation time increased. This result indicates that the energy needed to form CTCs was decreased by LPUV irradiation. LPUV irradiation depends on the chemical structure and polymer cohesiveness. This result is consistent with those shown in Fig. 9.

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Fig. 14 Fluorescence spectra of 6FDA-TeMPD films LPUV-irradiated from 0 to 10 J/cm2 at an excitation wavelength of 325 nm. (Reprinted with permission from J. Photopolym. Sci. Technol. 25, 401. (Copyright (2012) The Society of Photopolymer Science and Technology))

Figure 14 shows that the maximum peak emission intensities of all 6FDATeMPD films increased as the LPUV irradiation time increased. Similar to those of the UV spectra, the peak emission intensity ratio of all prepared films can be expressed as follows: Peak emission intensity raito ¼

I yJ  I 0J I 0J

(2)

where IyJ is the peak emission intensity for y J/cm2 LPUV irradiation and I0J is the peak emission intensity without LPUV irradiation. The peak emission intensity ratios of all films as a function of the LPUV irradiation time are presented in Fig. 15. As the irradiation time increased, the peak emission intensity of the 6FDA-TeMPD films thermally treated at 250
C (THF 250
C and DMAc 250
C) increased. On the other hand, a time lag in the increase in the 0–2 J/cm2 range was observed in the 6FDA-TeMPD films thermally treated at 100
C (THF 100
C and DMAc 100
C). The order of the increase is DMAc 250
C  THF 250
C > THF 100
C  DMAc 100
C. LPUV irradiation for the photoreaction strongly depended on polymer cohesiveness. Figures 13b and 15 show that the photoreactivity of the THF 250
C 6FDA-TeMPD film was the highest among all prepared films, whereas that of the DMAc 100
C 6FDATeMPD film was the lowest. The photoreactivity and polymer cohesiveness of the films varied with the film preparation protocol. Therefore, the photoalignments of the THF 250
C and DMAc 100
C 6FDA-TeMPD films, which exhibited the highest and lowest photoreactivities, respectively, were systematically investigated.

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Fig. 15 Peak emission intensity ratio of the maximum peak in the fluorescence spectra as a function of LPUV exposure from 0 to 10 J/cm2. Film preparation conditions: THF 100
C (○), THF 250
C (●), DMAc 100
C (□), DMAc 250
C (■). (Reprinted with permission from J. Photopolym. Sci. Technol. 25, 401. (Copyright (2012) The Society of Photopolymer Science and Technology))

Fig. 16 POM images of LPUV-irradiated 6FDA-TeMPD films from 0 to 10 J/cm2. (Reprinted with permission from J. Photopolym. Sci. Technol. 25, 401. (Copyright (2012) The Society of Photopolymer Science and Technology))

The POM images of the LC cells in the 6FDA-TeMPD films are shown in Fig. 16. The images with the polarizer parallel (0
) to the LPUV irradiation direction were obtained with crossed Nicols, whereas those with the polarizer perpendicular (90
) to the LPUV irradiation direction were obtained with parallel Nicols. The 6FDA-

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TeMPD films prepared without LPUV irradiation showed no photoalignment characteristics, whereas the 6FDA-TeMPD films irradiated with LPUV above 1 J/cm2 exhibited distinct photoalignment. The LC molecules were oriented in the direction of the LPUV irradiation. However, no significant difference was observed between the photoalignments of the films with the highest and lowest photoreactivities under LPUV irradiation higher than 1 J/cm2. Figures 14 and 15 show that the polymer cohesiveness of the DMAc 100
C 6FDA-TeMPD film decreased at 1 J/cm2 LPUV irradiation. However, this film did show photoalignment characteristics. Therefore, the photoalignment characteristics were determined by the chemical structure and did not depend on polymer cohesiveness. No difference was observed between the photoalignment characteristics of the different photoreactive films. Moreover, the photoalignment sensitivity was determined by comparing the two 6FDA-TeMPD films irradiated at 0.5 J/cm2 LPUV. The THF 250
C 6FDA-TeMPD film exhibited photoalignment characteristics. However, the photoalignment exhibited by the DMAc 100
C 6FDA-TeMPD film was low. The photoalignment sensitivity increased as the photoreactivity increased. In addition, the photoalignment sensitivity did not depend on polymer cohesiveness but rather on the residual solvent trace in the film. The 6FDA-TeMPD film prepared under low volatility preparation conditions exhibited photoalignment characteristics because the residual solvent disrupted the cross-linking polymer segments through photoreaction. This result indicates that photoalignment sensitivity can be increased by reducing the residual solvent in the polymer film.

Photoalignment Characteristics of CT Complex-Type Polyimides with Tetramethyl Phenylene Diamine In the 6FDA-TeMPD polyimide, the diamine component, TeMPD, generally has a high electron-releasing property, whereas the acid component, 6FDA, has an electron-accepting property. A CT transition can take place because of the transition from TeMPD to the 6FDA site. In general, the behavior is known to occur because of the formation of CTCs or CT interaction in polyimides. The CTCs are formed between the polymer chains in 6FDA-TeMPD polyimide. Based on our previous study, we expected that the CT interaction between polymer segments strongly depended on the photoreaction upon LPUV irradiation. However, no direct evidence for these polyimides formed CTCs has been presented in the literature. The energies of the LUMO, eLUMO, of the acid component and those of the HOMO, eHOMO, of the diamine component depend on CTC formation (Choi et al. 2009). The eLUMO and eHOMO are measures of the electron-accepting and electron-releasing properties observed during CTC formation, respectively. As the eLUMO decreases and the eHOMO increases, it is easy to form CTCs because of the lower transition energy between HOMO and LUMO. Based on the MO calculation, it can be concluded that TeMPD shows high eHOMO in 6FDA-TeMPD polyimides (8.71 eV). In the current study, the photoalignment characteristics of TeMPD-based polyimides with various eLUMO of the acid component were systematically investigated.

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Because the transition energies between HOMO and LUMO were different, the CT energies were different. The chemical structures of these TeMPD-based polyimides are shown in Fig. 17. 3,3,4,4-Diphenylsulfonetetracarboxylic dianhydride (DSDA) (eLUMO = 2.22 eV), 6FDA (eLUMO = 1.98 eV), and cyclobutane-1,2,3,4-tetracarboxylic dianhydride (CBDA) (eLUMO = 0.50 eV) were used in the current study. The effect of transition energy between eHOMO and eLUMO on photoalignment and reactivity were also analyzed in the DSDA-TeMPD, 6FDA-TeMPD, and CBDATeMPD polyimides. The chemical structure of each product, as shown in Fig. 18, was confirmed by FTIR analysis. The isotropic, dense polyamide films were prepared by spin-coating. These three polyimide films were obtained using a thermal imidization method.

Fig. 17 Chemical structures and frontier MOs of TeMPD, DSDA, 6FDA, and CBDA. (Reprinted with permission from J. Photopolym. Sci. Technol. 26, 769. (Copyright (2013) The Society of Photopolymer Science and Technology))

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Fig. 18 Thermal imidization reactions of TeMPD-based polyimides. (Reprinted with permission from J. Photopolym. Sci. Technol. 26, 769. (Copyright (2013) The Society of Photopolymer Science and Technology))

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The FTIR spectra of LPUV-irradiated TeMPD-based polyimides for 10 J/cm2, unirradiated TeMPD-based polyimides (after thermal imidization), and TeMPDbased polyamides are shown in Fig. 19. As shown in Fig. 19a, the N–H stretching bond peak was observed near 1650 cm1 in DSDA-TeMPD polyamide. However, the peak was not observed after thermal imidization. In contrast, the intensity of the O–H stretching peak near 3300 cm1 decreased and that of the C=O stretching peak near 1750 cm1 increased after thermal imidization. In the case of 6FDA-TeMPD, the N-H bond peak derived from polyamide was observed before thermal imidization as shown in Fig. 19b. For DSDA-TeMPD also, the N-H bond peak was not observed after thermal imidization. In addition, the O–H stretching peak intensity increased and the C=O stretching peak intensity decreased upon LPUV irradiation. These results are consistent with those in our previous study (Sato et al. 2011). The N-H bond peak was observed for CBDATeMPD after thermal imidization as shown in Fig. 19c. The peak intensity slightly decreased after thermal imidization. CBDA-TeMPD was not fully imidized. Through LPUV irradiation, the O–H stretching and N-H bond peaks intensities increased. In contrast, the C=O stretching peak intensity slightly decreased. These results indicate that the photoreaction mechanism of DSDATeMPD is similar to that of 6FDA-TeMPD. Therefore, an intermolecular crosslinked photoreaction owing to CT interaction occurred in DSDA-TeMPD, as shown in our previous study (Sato et al. 2011, 2012). In contrast, the intermolecular cross-linked photoreaction only partially occurred in CBDATeMPD polyimide. The thermal imidization ratio can be expressed by comparing the O–H stretching peak’s intensity, I, as follows: (I(before thermal imidization) - I(after thermal imidization))/ I(before thermal imidization). The ratios of DSDA-TeMPD, 6FDA-TeMPD, and CBDA-TeMPD were 99.7%, 96.2%, and 63.4%, respectively. CBDA-TeMPD exhibited a lower reaction ratio than DSDA-TeMPD and 6FDA-TeMPD because of low main-chain rigidity. The thermal imidization reaction of CBDA-TeMPD proceeds weakly. The CBDA-TeMPD polyimide after thermal imidization contained approximately half polyamide. The polyamide depends on the photoalignment characteristics. It is expected that the photoalignment characteristic of CBDA-TeMPD is different from those of DSDA-TeMPD and 6FDA-TeMPD. The UV spectra from 2.5 to 6.5 eV of these LPUV-irradiated TeMPD-based polyimides for 10 J/cm2, unirradiated TeMPD-based polyimides, and TeMPDbased polyamides are shown in Fig. 20. Strong peaks near 6.2 eV were observed in DSDA-TeMPD and 6FDA-TeMPD. The peak position shifted to 5.5 eV after thermal imidization, and the energy after imidization was lower than that before imidization. In addition, the peak intensity decreased after LPUV irradiation. Moreover, weak peaks near 4.4 eV were observed in DSDA-TeMPD and 6FDA-TeMPD before imidization. The peak position shifted to 4.0 eV after thermal imidization and LPUV irradiation. In contrast, a strong peak near 6.0 eV was observed in CBDATeMPD before thermal imidization; the peak intensity decreased without a change in the peak position. A weak peak near 4.5 eV was newly found after thermal imidization and LPUV irradiation. There was no change in the intensity and position

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Fig. 19 IR absorption spectra of LPUV-irradiated and unirradiated TeMPD-based polyimides and polyamides. Polyimide: (a) DSDA-TeMPD, (b) 6FDA-TeMPD, (c) CBDA-TeMPD. (Reprinted with permission from J. Photopolym. Sci. Technol. 26, 769. (Copyright (2013) The Society of Photopolymer Science and Technology))

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Fig. 20 Absorption spectra of LPUV-irradiated and unirradiated TeMPD-based polyimides and polyamides. (Reprinted with permission from J. Photopolym. Sci. Technol. 26, 769. (Copyright (2013) The Society of Photopolymer Science and Technology))

of the weak peak. The peaks changed by thermal imidization and LPUV irradiation were assigned based on the results of the MO calculation. For the TeMPD-based polyamides and polyimides, orbital compositions of the frontier MOs are shown in Fig. 21. They are mainly described by one electron excitation from the HOMO to the LUMO. Consequently, the HOMO!LUMO transition implies electron density transfer. Note that the charge at HOMO and LUMO is located on the TeMPD and acid component sites, respectively. That is, photoabsorption of polyamides and polyimides in the current study depends on the CT transition from the TeMPD site to the acid component site because of the phenyl group π–π transitions. Calculated UV absorptions of these TeMPD-based polyamides and polyimides are shown in Fig. 22. The peak position slightly shifted toward 5.3 eV for polyamides and 5.0 eV for polyimides in the case of DSDATeMPD and 6FDA-TeMPD. The peak shifts depend on the imidization reaction, as shown in Fig. 20. On the other hand, there was no change in the peak position between polyamide and polyimide in CBDA-TeMPD. This result is in good agreement with the observed UV spectra in Fig. 20. This behavior is well-known with regard to the formation of CTCs or CT interactions in polyimides (Kanehashi et al. 2011). This interaction is attributed to the presence of π electrons between the ring structures in polyimides. The CTC structure is similar to a sandwich structure between imides and the neighboring benzene ring. The CT interaction observed during CTC production strongly depends on the fluorescence spectrum. Figure 23 shows the fluorescence spectra of TeMPDbased polyamides and polyimides before and after thermal imidization and after LPUV irradiation. These polymers exhibited an emission spectrum band at 470 nm after excitation at 325 nm. This band is attributed to the intramolecular interactions

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Fig. 21 HOMO and LUMO compositions of frontier MOs for TeMPD-based polyimides and polyamides. (Reprinted with permission from J. Photopolym. Sci. Technol. 26, 769. (Copyright (2013) The Society of Photopolymer Science and Technology))

of the aromatic polyimide that contains an alternating sequence of electron-rich donor and electron-deficient acceptor molecules. Therefore, the CTC formation and polymer cohesiveness can be determined via the emission peak intensity. The maximum peak emission wavelength was approximately 470 nm. Figure 23 shows that the maximum peak emission intensities of all TeMPD-based polyimides increased after LPUV irradiation. The LPUV irradiation depends on the chemical structure and polymer cohesiveness. It is expected that the CT interaction between polymer segments strongly depends on the photoreaction. The intermolecular crosslinked photoreaction derived from CT interaction occurred, as shown in our previous study (Sato et al. 2011a, 2012). The increases in the peak emission intensity of DSDATeMPD and 6FDA-TeMPD were more remarkable than that in CBDA-TeMPD. This is because the transition energy from eHOMO to eLUMO of DSDA-TeMPD and 6FDATeMPD was lower than that of CBDA-TeMPD, as determined by MO calculations. As shown in Fig. 21, the values of DSDA-TeMPD, 6FDA-TeMPD, and CBDA-TeMPD

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Fig. 22 Calculated absorption spectra of TeMPDbased polyimides and polyamides. (Reprinted with permission from J. Photopolym. Sci. Technol. 26, 769. (Copyright (2013) The Society of Photopolymer Science and Technology))

Fig. 23 Fluorescence spectra of LPUV-irradiated and unirradiated TeMPD-based polyimides and polyamides excited at 325 nm. (Reprinted with permission from J. Photopolym. Sci. Technol. 26, 769. (Copyright (2013) The Society of Photopolymer Science and Technology))

polyimides were 7.398 eV (9.248 eV ! 1.850 eV), 7.404 eV (9.211 eV ! 1.807 eV), and 9.098 eV (9.209 eV ! 0.111 eV), respectively. These results indicate that the intermolecular cross-link type photoreactivity can be increased by reducing the transition energy from eHOMO to eLUMO.

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Fig. 24 POM images of LPUV-irradiated and unirradiated TeMPD-based polyimides. (Reprinted with permission from J. Photopolym. Sci. Technol. 26, 769. (Copyright (2013) The Society of Photopolymer Science and Technology))

The photoalignment of the LC molecules in TeMPD-based polyimides was investigated. The POM images of the LC cells in the TeMPD-based polyimide films are shown in Fig. 24. All TeMPD-based polyimide films prepared without LPUV irradiation showed no photoalignment characteristics, whereas those irradiated with 10 J/cm2 LPUV exhibited distinct photoalignment. All the color variation domains of the POM image were the same. The LC molecules were oriented in the same direction. However, the photoalignment of the CBDA-TeMPD polyimide film was similar to that of DSDA-TeMPD and 6FDA-TeMPD, in contrast to the IR result. The imidization ratio and the photoreactivity of polyimide did not depend on the photoalignment characteristic. The photoalignment was only dependent on surface photoreactivity. The chemical structure of polyimide did not affect surface photoreactivity and the apparent photoalignment.

Conclusions We have summarized new photoalignment polyimides that are different from the existing photoalignment polymers such as azo- and coumarin-containing polymers. Then, we theorized that the CT interaction between the polymer segments strongly depended on the photoreaction caused by LPUV irradiation. Polymers with photoalignment characteristics will be required soon. The development of the new polymers will also be conducted while maintaining high heat resistance, high solubility in polar solvents, and high optical transparency based on

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our design method. We can expect that the polymers would extend to further application fields such as holograms and resist materials.

References Choi M-C, Wakita J, Ha C-S, Ando S (2009) Highly transparent and refractive polyimides with controlled molecular structure by chlorine side groups. Macromolecules 42:5112–5120 Cognard J (1982) Alignment of Nematic liquid crystals and their mixtures. Gordon and Breach Science, London Gibbons WM, Shannon PJ, Sun ST, Swetlin BJ (1991) Surface-mediated alignment of nematic liquid-crystals with polarized laser-light. Nature 351:49–50 Ha KR, West JL (2002) Studies on the photodegradation of polarized UV-exposed PMDA-ODA polyimide films. J Appl Polym Sci 86:3072–3077 Kanehashi S, Nakagawa T, Nagai K, Duthie X, Kentish S, Stevens G (2007) Effects of carbon dioxide-induced plasticization on the gas trans-port properties of glassy polyimide membranes. J Membr Sci 298:147–155 Kanehashi S, Sato S, Nagai K (2011) Membrane color and gas permeability of 6FDA-TeMPD polyimide membranes prepared with various membrane preparation protocols. Polym Eng Sci 51:2360–2369 Miyata S, Sato S, Nagai K, Nakagawa T, Kudo K (2008) Relationship between gas transport properties and fractional free volume determined from dielectric constant in polyimide films containing the hexafluoroisopropylidene group. J Appl Polym Sci 107:3933–3944 Nishikawa M, West JL (1999) Order parameter of liquid crystal on polyimide with polarized ultraviolet-light exposure. Jpn J Appl Phys 38:L331–L333 Sato S, Suzuki M, Kanehashi S, Nagai K (2010) Permeability, diffusivity, and solubility of benzene vapor and water vapor in high free volume silicon- or fluorine-containing polymer membranes. J Membr Sci 360:352–362 Sato S, Ito H, Mizunuma T, Nagai K, Matsumoto H, Matsumoto S (2011a) Substitute effect of fluorine-containing polyimides with hexafluoroisopropylidene group on photo alignment of liquid crystal molecule. J Photopolym Sci Technol 24:617–623 Sato S, Ose T, Miyata S, Kanehashi S, Ito H, Matsumoto S, Iwai Y, Matsumoto H, Nagai K (2011b) Relationship between the gas transport properties and the refractive index in high-free-volume fluorine-containing polyimide membranes. J Appl Polym Sci 121:2794–2803 Sato S, Gondo D, Sugiyama H, Nagai K, Matsumoto S, Matsumoto H (2012) Photo-reactivities of 6FDA-TeMPD polyimides prepared using different film preparation protocols as an alignment layer. J Photopolym Sci Technol 25:401–407 Sato S, Ichikawa M, Ose T, Miyata S, Takahashi Y, Kanehashi S, Matsumoto H, Nagai K (2013) Gas transport and optical properties of ABA-type triblock copolymers designed using fluorinecontaining polyimide macroinitiators with methyl methacrylate. Polym Int 62:1377–1385 Schadt M, Schmitt K, Kozinkov V, Chigrinov V (1992) Surface-induced parallel alignment of liquid-crystals by linearly polymerized photopolymers. Jpn J Appl Phys 31:2155–2164 Shitomi H, Ibuki T, Matsumoto S, Onuki H (1999) Optically controlled alignment of liquid crystal on polyimide films exposed to undulator radiation. Jpn J Appl Phys 38:176–179

Gas Permeation and Barrier Properties of Liquid Crystalline Polymers

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Shota Ando, Shuichi Sato, and Kazukiyo Nagai

Contents Definition . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . Introduction . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . Gas Permeation Mechanism in Crystalline and Liquid Crystalline Polymers . . . . . . . . . . . . . . . . . Statistical Literature Data Analysis in Crystalline and Liquid Crystalline Polymers . . . . . . . . . . Gas Permeability in Crystalline and Liquid Crystalline Polymers . . . . . . . . . . . . . . . . . . . . . . . . . . . . . Gas Solubility in Crystalline and Liquid Crystalline Polymers . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . Gas Diffusivity in Crystalline and Liquid Crystalline Polymers . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . Analysis of the Relationship Between Gas Transport Properties and Crystalline Structure in Crystalline and Liquid Crystalline Polymers . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . Relationship Between Experimentation and Extrapolation in Crystalline and Liquid Crystalline Polymers . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . Relationship of Gas Permeability, Solubility, and Diffusivity . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . References . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . .

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Abstract

This chapter reviews gas permeation and barrier properties of crystalline and liquid crystalline polymers from both theoretical mechanism and experimental measurement. The permeability is closely related to both solubility and diffusivity in polymers. Various factors such as crystallinity and crystalline-amorphous morphology play important roles in the solubility and diffusivity, and thus the permeability of crystalline and liquid crystalline polymers.

S. Ando · K. Nagai (*) Department of Applied Chemistry, Meiji University, Kawasaki, Kanagawa, Japan e-mail: [email protected]; [email protected] S. Sato Department of Electrical and Electronic Engineering, Tokyo Denki University, Adachi-ku, Tokyo, Japan © Springer Nature Switzerland AG 2020 L. Zhu, C. Y. Li (eds.), Liquid Crystalline Polymers, Polymers and Polymeric Composites: A Reference Series, https://doi.org/10.1007/978-3-030-43350-5_67

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Keywords

Crystalline · Gas transport property · Liquid polymer · Crystalline structure · Gas barrier · Phase transition

Definition The gas transport properties of a polymer membrane are not only affected by its physical properties but also by the difference of its crystalline structure (e.g., crystalline, liquid crystalline, and amorphous). Several studies have reported the relationship between crystalline structure and gas transport properties, and this relationship is important in the design of various polymer materials. In this chapter, the effects of permeability, solubility, diffusivity of each polymer film on the transport properties of various gas molecules are explained based on data from literature.

Introduction Recently, polymer science research has focused on membrane materials that have precise, highly ordered structures, such as those that are crystalline or have a wellordered orientation. Polymers are generally amorphous or semi-crystalline. The latter is also called a partially crystalline polymer and consists of a complex combination of crystalline and amorphous regions. In the crystalline region, polymer segments are packed parallel to each other and form a phase of finite size. Small molecules only pass through the amorphous region because the polymer chain mobility is considerably restricted in the crystalline phase, thereby causing difficulties in diffusion and dissolution (Paul and Yampol’ skii 1994). To date, a large number of physical and transport properties for different gases, including oxygen (O2), nitrogen (N2), and carbon dioxide (CO2) have been studied in crystalline and liquid crystalline polymers. As shown in Fig. 1, these polymers include polyethylene (PE) (Waack et al. 1955; Alter 1962; Michaels and Parker 1959; Michaels and Bixler 1961a; Bixler et al. 1963; Holden et al. 1985; Bernal-Lara et al. 2005; Mendes et al. 2004; Compan et al. 1998, 1997; Kofinas et al. 1994; Ash et al. 1970; Srinivas et al. 2003), poly[bis(trifluoroethoxy phosphazene)] (PTFEP) (Hirose et al. 1989; Nagai et al. 2000; Mizoguchi et al. 1991), polypropylene (PP) (Kofinas et al. 1994; Vieth and Wuerth 1969; Somlai et al. 2005; Lin et al. 2008a, b; Incarnato et al. 2000), poly(4-methyl-1-pentene) (PMP) (Puleo et al. 1989; Mohr and Paul 1991; Levaesalmi and McCarthy 1995), poly(lactic acid) (PLA) (Komatsuka et al. 2008; Komatsuka and Nagai 2009; Sawada et al. 2010; Lehermeier et al. 2001; Auras et al. 2003; Bao et al. 2006), polyethylene terephthalate (PET) (Michaels et al. 1963; Qureshi et al. 2000; Polyakova et al. 2001; Hu et al. 2002a, 2005a, b; Liu et al. 2004), polystyrene (PS) (Hodge et al. 2001; Prodpran et al. 2002), polyacrylate (PA) (Mogri and Paul 2000a, b, 2001a, b; O’Leary and Paul 2006a, b;

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Fig. 1 Chemical structure of each polymer

Kirkland and Paul 2008), polyethylene naphthalate (PEN) (Hu et al. 2002b), poly (ethylene oxide) (PEO) (Lin and Freeman 2004; Pethe et al. 2008; Metz et al. 2004), poly(vinylidene fluoride) (PVDF) (El-Hibri and Paul 1986), liquid crystalline polyester (Incarnato et al. 2000; Chiou and Paul 1987; Weinkauf and Paul 1991, 1992a, b; Weinkauf et al. 1992; Park et al. 1996; Flodberg et al. 2003; De Candia et al. 1990, 1991), polyamide (Waack et al. 1955; Ash et al. 1970; Hu et al. 2005c; Kanekura et al. 2005), among others (crystalline (Tsujita et al. 1990; Kumazawa et al. 1994; Uriarte et al. 1998; Wang and Easteal 1999; Kim and Lee 2001; Kim et al. 2001; Mensitieri et al. 1996; Laguna et al. 2003; Buquet et al. 2009; Cowling and Park 1979) and liquid crystalline polymers (Kajiyama et al. 1982, 1985, 1988; Chen et al. 1991, 1992; Chen and Hsiue 1993; Kawakami et al. 1997; Hu et al. 2003, 2006a, b, 2007)). Other studies have also found that a low-density PE (density: 0.910–0.925 g cm
3) has a crystallinity of approximately 60%, whereas high-density PE (density: 0.94–0.97 g cm
3) has a crystallinity of approximately 90% (Mark 2007). This result suggested that the density of crystalline polymers is directly related to the degree of crystallinity. Furthermore, earlier studies showed that O2 permeability decreases with increasing density in PE (Alter 1962). This result can be attributed to the decrease in both gas solubility and diffusivity. However, we recently reported that gas transport properties were independent of PLA crystallinity (Komatsuka et al. 2008; Komatsuka and Nagai 2009; Sawada et al. 2010). This result suggests that gas transport properties are strongly affected by the size and distribution of the crystalline phase rather than by the polymer crystallinity (Komatsuka et al. 2008). Furthermore, unlike common crystalline polymer membranes, the permeability of crystalline PLA membranes is higher

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than that of amorphous PLA membranes. One possible explanation based on this result is that a continuous space for gas diffusion may exist around the interface between the crystalline and amorphous regions (Sawada et al. 2010). In the same way, permeability in PS is also independent of its crystallinity (Hodge et al. 2001; Prodpran et al. 2002). Syndiotactic PS (s-PS) is known to have several crystalline structures (i.e., α, β, γ, and δ crystalline). These structures not only affect the physical properties of the polymer but their gas transport properties as well (Hodge et al. 2001; Prodpran et al. 2002). This behavior is similar to PLA membranes. In general, the gas permeability of the β-form crystalline structure is lower than that of the α-form because the β-form structure has more tightly packed polymer chain segments. The density of the β-form is higher than that of the amorphous polymers, whereas the density of the α-form is lower than that of the amorphous polymers. The gas permeability and solubility of the β-form decrease with increasing crystallinity, and its diffusivity remains constant with increasing crystallinity. The gas permeability and diffusivity of the α-form structure increase and the solubility decreases with increasing crystallinity. Based on these results, porous microcrystalline structures with nanochannels may be adopted to increase gas diffusivity, thereby increasing gas permeability (Hodge et al. 2001; Prodpran et al. 2002). Furthermore, s-PS with molecular cavities of different shapes and sizes were obtained from an α-form crystal (Tsujita et al. 2005). Therefore, each crystalline polymer has an intrinsic crystalline phase with different sizes and distributions. These phases affect not only their physical and thermal properties but also their gas transport properties. The density of crystalline PMP is slightly lower than that of its amorphous counterpart (Griffith and Ranby 1960; Ranby et al. 1962). This result was considered based on the characteristic crystalline structure of PMP. The bulky pendant structure in PMP inhibits close polymer segments from packing and results in a low-density crystal with gaps of almost 4 Å between chain segments (Puleo et al. 1989). Therefore, a small molecule can pass though this characteristic crystalline structure in PMP. By comparison, the properties of liquid crystalline polymers are different from those of amorphous and crystalline polymers in terms of fluidity and orientation. These kinds of polymers have three types of phases, namely, nematic, cholesteric, and smectic phases. As shown in Fig. 1, glassy liquid crystalline polymers, such as poly( p-hydroxybenzoic acid-co-6-hydroxy-2-naphthioic acid) (HBA/HNA) and poly( p-phenylene terephthalamide) (PPTA), show remarkably high gas barrier properties. For example, the O2 permeabilities of HBA/HNA copolymer and PPTA are lower than that of PET or PVC but are similar to that of polyacrylonitrile (Weinkauf and Paul 1991; Weinkauf et al. 1992). These high gas barrier properties are attributed to the efficient packing of the molecules in the solid state. The gas barrier property of the HBA/HNA copolymer also increases with increasing HNA content. The napthyl unit restricts the rotation mobility of the main chain and packed polymer chain segments; thus, high gas barrier properties were developed. Interestingly, in spite of the low crystallinity (13–26%) of the HBA/HNA copolymer, it has a high gas barrier property (Weinkauf and Paul 1992b). However, the solubility of the liquid crystalline polymer in organic solvents, such as acetone, is quite low (Freeman and Hill 1998).

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As mentioned above, the microstructure of crystalline and liquid crystalline polymers strongly affect not only their physical properties but also their transport properties. We have already reported and discussed the validity of dual mode sorption parameters and gas transport properties based on a statistical analysis employing 250 types of glassy amorphous polymers from an exhaustive search of published data (Kanehashi and Nagai 2005). In the present study, published data on the physical and gas transport properties of 300 crystalline and liquid crystalline polymers were statistically evaluated and analyzed in comparison with glassy amorphous polymers (Kanehashi and Nagai 2005).

Gas Permeation Mechanism in Crystalline and Liquid Crystalline Polymers In general, gas permeation (P) in polymeric dense membranes is explained through solution/diffusion mechanisms (Paul and Yampol’ skii 1994). P ¼SD

(1)

where S is the solubility and D is the diffusivity. Gas permeation in partially crystalline and liquid crystalline polymers is illustrated through a dual phase model because these polymers have a combination of amorphous and crystalline regions (Paul and Yampol’ skii 1994). Solubility, S, in crystalline polymers is described by the following equation: S¼

  100
X c Sa 100

(2)

where Xc is the degree of crystallinity and Sa is the solubility for a purely amorphous polymer. Earlier studies showed that the crystalline PE solubility can be described in this equation (Michaels and Parker 1959; Michaels and Bixler 1961). However, the solubility for amorphous PE is two times higher than that expected for this model (Budzien et al. 1998). These results suggest that the crystalline phase may have affected the gas solubility and diffusivity of the amorphous phase. The simple relation of the crystalline phase on diffusivity is given by D¼

  100
Xc Da 100

(3)

where Da is the diffusivity of the purely amorphous polymer (Michaels and Bixler 1961; Vieth and Wuerth 1969; Michaels et al. 1963; Lasoski and Cobbs 1959). The correlation of diffusivity with crystallinity is empirical, and several studies have reported that the exponent ranges from 1 to 2 (Michaels and Bixler 1961; Vieth and Wuerth 1969). Equation 3 seems to be a rough explanation for gas diffusivity in crystalline and liquid crystalline polymers and can only estimate the effect of

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crystallinity. This equation is inadequate in explaining the diffusivity in crystalline and liquid crystalline polymers because the crystalline phase seems to include two factors (Michaels and Bixler 1961). Small molecules are assumed to be incapable of diffusing or permeating through the crystalline phase. Therefore, they must pass through an elongated path because of the presence of the crystalline phase. This factor is generally accounted for as the tortuosity or geometric impedance factor, τ. The other factor is the chain-immobilization factor, β. The reciprocal of β, (β
1), represents a fractional reduction in diffusivity because of the restriction of the polymer chain segmental mobility in the amorphous phase. The effect of the crystalline phase on diffusivity can be expressed using τ and β. D¼

Da τβ

(4)

The size of the crystalline phase is known to increase with increasing τ, and the molecular size is known to increase with increasing β in several different crystalline PE (Michaels and Bixler 1961). τ is also correlated with the amorphous fraction, α, in PET (Michaels et al. 1963). τ¼

1 α

(5)

This equation also applies to vapor diffusion in PET (Lasoski and Cobbs 1959). In addition, Eq. 5 for PE can also be written as follows. τ ¼ α
n

(6)

where n varies from 1.25 to 1.88 for crystalline PE (Michaels and Bixler 1961). Although the τ and β of particular crystalline polymers, such as PE (Michaels and Parker 1959; Michaels and Bixler 1961; Bixler et al. 1963), PET (Michaels et al. 1963), and PEO (Metz et al. 2004), have been reported, discussions comparing these parameters are quite few. Based on this model, the gas permeability in a crystalline polymer can be written using Eqs. 2 and 4.  P¼

   100
X c Da 100
X c ¼ Sa  S a  γDa 100 τβ 100

γ¼

1 P ¼ τβ ð100
X c =100ÞS a  Da

(7)

(8)

where γ is the diffusion contribution parameters in crystalline and liquid crystalline polymers. If γ > 1, the crystalline phase may enhance gas diffusion, that is Da < D in Eq. 4. If γ < 1, the crystalline phase may restrict gas diffusion, that is, Da > D.

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Statistical Literature Data Analysis in Crystalline and Liquid Crystalline Polymers Statistical analyses were performed based on approximately 300 different polymers (250 crystalline and 50 liquid crystalline). This study employed the published gas transport properties determined at 25–35  C with 1–10 atm of feed pressure for crystalline and those determined at 20–40  C and 1–10 atm for liquid crystalline polymers. The crystallinity, glass transition temperature (Tg), and maximum crystallinity [Xc(max)] of the polymers used in this study are summarized in Table 1. Polymers are defined as a glassy state below the glass transition temperature and a rubbery state above the glass transition temperature (Mark 2003). The polymers that are rubbery at the measurement temperature were (Xc(max)%): PE (93%) (Maxwell et al. 1996), PTFEP (>60%) (Mizoguchi et al. 1991), PP (89%) (Farrow 1963), and PMP (96%) (Suzuki et al. 2002), and the polymers that are glassy state at the measurement temperature were (Xc(max)%): PLA (>70%) (Renouf-Glauser et al. 2005), PET (90%) (Köncke et al. 1996), PS (>60%) (Wu et al. 2001), PA (>50%) (Mogri and Paul 2001a), PEN (87%) (Li et al. 2001), HBA/HNA (>40%) (Wiberg and Gedde 1997), and PPTA (81%) (Hindeleh and Abdo 1989), where the latter two are liquid crystalline polymers. The maximum degrees of crystallinity of the crystalline and liquid crystalline polymers were determined using SciFinder® (CAS a division of the American Chemical Society). The maximum crystallinity of many polymer samples in this study was over 80% with the exception of PTFEP, PLA, PS, PA, and HBA/HNA. The Tg values were also obtained from literature (Krevelen and Nijenhuis 2009). This study also employed previously reported data on glassy amorphous polymers for comparison (Kanehashi and Nagai 2005). The O2 transport properties in crystalline and liquid crystalline Table 1 Crystalline and liquid crystalline polymers used in this study

Polymer PE PTFEP PP PMP PLA PET PS PA PEN HBA/HNA (LCP) PPTA (LCP)

Type Crystalline Crystalline Crystalline Crystalline Crystalline Crystalline Crystalline Crystalline Crystalline Liquid crystalline Liquid crystalline

Crystallinity range (%) 22–78 31–60 38–62 22–66 4–40 1–51 2–37 24–50

State at experimental temperature Rubbery Rubbery Rubbery Rubbery/glassy Glassy Glassy Glassy Glassy

Maximum crystallinity (%) 95 >60 89 96 >70 90 >60 >50

Tg ( C)
78
70
10 30 60 69 100 104

1–39 13–26

89 >40

113 –

Glassy Glassy

22–45

81

–

Glassy

(a)

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polymers have often been investigated because these are crucial in electrical devices and packaging materials. In these applications, an O2 barrier property is desired. Data for other gases, such as CO2 and hydrocarbons, are few. This study focused on the gas transport properties of CO2 as a condensable gas, N2 as a non-condensable gas, and O2 as an often reported gas in crystalline polymers. Data for gas transport properties were determined via a differential or equal pressure method. The degree of crystallinity was determined via several methods, including X-ray diffraction (XRD), thermal analysis, vibrational spectroscopy, volumetric method, and solidstate nuclear magnetic resonance (NMR) (Runt and Kanchanasopa 2008). The crystallinity calculated via XRD, density, and differential scanning calorimetry was used in this analysis.

Gas Permeability in Crystalline and Liquid Crystalline Polymers Figure 2 shows the gas permeability values of crystalline and liquid crystalline polymers as a function of crystallinity. Liquid crystalline polymers have higher gas barrier properties (i.e., lower gas permeability) compared with crystalline polymers. Gas permeability is known to decrease with increasing crystallinity in crystalline and liquid crystalline polymers because gases cannot diffuse and permeate in the crystalline phase. However, this study shows that gas permeability in the crystalline and liquid crystalline polymers are independent of their crystallinity at lower crystallinity values and are almost the same or higher than those of amorphous polymers. For example, the O2 permeability of crystalline and liquid crystalline polymers is almost constant from 20% to 60% crystallinity for PMP, 10% to 30% for liquid crystalline HBA/HNA, and 20% to 40% for PPTA. By contrast, the permeabilities of PMP, PLA, and PPTA decrease from 60% to 65%; 25% to 40%; and 40% to 50%, respectively. This reduction is also observed in PP from 40% to 60%, in PA from 20% to 50%, in PET from 0% to 50%, and in PEN from 0% to 40%. Interestingly, the O2 permeability slightly increases between 0% and 25% for PLA, and 0% and 40% crystallinity for PS. Only the O2 permeability in PE decreases from 20% to 60% and is sharply reduced from 60% to 80%. The reduction in PE was larger than those in the other crystalline and liquid crystalline polymers. As shown in PMP and PLA, the gas permeability begins to decrease beyond a specific crystallinity value. This result suggests that the phase transitions of the crystalline structure in crystalline and liquid crystalline polymers occur at these crystallinity values. Hence, the gas permeability of crystalline and liquid crystalline polymers may decrease during phase transitions. The same pattern is observed for CO2 and N2 permeability.

Gas Solubility in Crystalline and Liquid Crystalline Polymers Figure 3 shows the gas solubility in crystalline and liquid crystalline polymers as a function of the crystallinity. Liquid crystalline polymers have higher gas barrier properties (i.e., lower gas solubility) compared with other crystalline polymers

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Gas Permeation and Barrier Properties of Liquid Crystalline Polymers

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Fig. 2 Gas permeability for (a) CO2, (b) N2, and (c) O2 in crystalline and liquid crystalline polymers as a function of crystallinity, Xc. (Reprinted with permission from J Membr Sci., 365, 40. (Copyright (2010) Elsevier)

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Fig. 3 (continued)

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Fig. 3 Gas solubility for (a) CO2, (b) N2, and (c) O2 and purely amorphous solubility in crystalline and liquid crystalline polymers as a function of crystallinity, Xc. (Reprinted with permission from J Membr Sci., 365, 40. (Copyright (2010) Elsevier)

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(Fig. 3). As shown, the O2 solubility of crystalline and liquid crystalline polymers slightly decreases when the crystallinity was increased from 0% to 40% for PLA, 0% to 50% for PET, 0% to 40% for PS, and 0% to 40% for PEN. The O2 solubility in PE also decreases when the crystallinity was increased from 20% to 78%, similar to other crystalline polymers. For example, the O2 solubility at crystallinities between 22% and 78% varies from 0.638 cm3 to 0.014 cm3 (STP) (cm3 atm)
1, thereby indicating a 45-fold difference within the same crystallinity for PE. For liquid crystalline polymers, the O2 solubility in HBA/HNA is almost constant between 10% and 30% crystallinity, whereas the value decreases in PPTA between 20% and 40% crystallinity. The solubility in other liquid crystalline polymers remains constant but somewhat scattered between the crystallinity of 0% and 50%. The CO2 and N2 solubilities in other polymers are similar to that of O2. As seen in these results, the solubility of many crystalline polymers, such as PE, PLA, PET, PA, and PEN, tends to decrease with increasing crystallinity, whereas the liquid crystalline polymers show constant values in this analysis. No increase in the solubility is observed for crystalline and liquid crystalline polymers. Figure 3 also shows the purely amorphous gas solubility, Sa, of crystalline polymers as a function of their crystallinity. Interestingly, regardless of the gas used, Sa in almost all crystalline and liquid crystalline polymers, with the exception of PEN, are independent of their crystallinity. Only the crystalline PEN shows a slight increase in solubility between 0% and 40%. This result suggests that the solubility in crystalline and liquid crystalline polymers with regards to Sa and crystallinity is consistent with Eq. 2.

Gas Diffusivity in Crystalline and Liquid Crystalline Polymers Figure 4 shows the gas diffusivity in crystalline and liquid crystalline polymers as a function of crystallinity. Liquid crystalline polymers generally have a higher gas barrier property (i.e., lower gas diffusivity) compared with other crystalline polymers. The O2 diffusivity, including their deviations, remains almost constant in the range of 10–30% for liquid crystalline HBA/HNA and 20–40% for PPTA. The diffusivity in PLA and PS slightly increases at crystallinities between 0% and 25% and 0% and 40%, respectively. On the other hand, the diffusivity in PET and PEN slightly decreases at crystallinities between 0% and 50% and between 0% and 40%, respectively. The O2 solubility in PE varies from 1.37  10
6 cm2 s
1 for 20% crystallinity to 3.30  10
9 cm2 s
1 for 78% crystallinity, thereby indicating a 415-fold decrease. These trends are also observed in both CO2 and N2.

Analysis of the Relationship Between Gas Transport Properties and Crystalline Structure in Crystalline and Liquid Crystalline Polymers Although gas transport properties (i.e., permeability, solubility and diffusivity) in crystalline and liquid crystalline polymers decrease with increasing crystallinity, based on Figs. 2, 3, and 4, gas permeability and diffusivity are not affected by

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Gas Permeation and Barrier Properties of Liquid Crystalline Polymers

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crystallinity at lower crystallinity ranges. Particular polymers show an almost constant or slight increase in gas permeability and diffusivity at a lower crystallinity range. The decrease in gas transport properties (i.e., permeability, solubility, and diffusivity) is observed in rubbery PE and in glassy PET and PEN. Interestingly, these values for liquid crystalline polymers remain unchanged for the range of crystallinity in this analysis. Gas permeability and diffusivity in PS increase with increasing crystallinity, whereas the gas solubility decreases. Although gas permeability and diffusivity for PLA slightly increase and then begins to decrease, solubility decreases with increasing crystallinity. Based on these results, gas transport properties do not always decrease at lower crystallinity values, whereas gas transport properties decrease at higher crystallinity values. These results suggest that the crystalline structure can form discontinuously in crystalline and liquid crystalline polymersat lower crystallinities. When the crystallinity was increased, the crystal structure gradually grows similar to a phase transition from the discontinuous domains. One possibility is that there may be a continuous space for gas diffusion around the interface between the crystalline and amorphous regions, which may be created because of the stress of polymer chain arrangements around the interface regions. The gas diffusion process is enhanced because the space is larger than the size of the gas molecule. Hence, gas permeability increases along with this increase in gas diffusion at lower crystallinities. Furthermore, gas permeability and diffusivity begin to decrease after the phase transition of the crystal structures. At a higher crystallinity, the crystal structure may effectively prevent the diffusion of gas molecules. Therefore, the crystal structure can be changed from discontinuous to continuous crystal domains after phase transition. Clearly, gas transport properties can be strongly affected by crystal structure parameters, such as crystal size and distribution, for crystalline and liquid crystalline polymers.

Relationship Between Experimentation and Extrapolation in Crystalline and Liquid Crystalline Polymers The purely amorphous gas transport parameters, Pa, Da, and Sa, in crystalline polymers were extrapolated using the best fit lines (Xc ! 0) in Figs. 2, 3, and 4 (Table 2). Different polymers show large differences in the gas transport parameters in their purely amorphous form. These results indicate that gas transport properties strongly depend on the chemical structure of the polymer. According to the solutiondiffusion model in polymeric dense membranes, CO2 and O2 permeability, Pa, in this analysis was consistent with solubility, Sa, and diffusivity, Da. By contrast, N2 permeability, Pa, and the product of the solubility, Sa, and diffusivity, Da, have large deviations based on the limited experimental data. The order of the gas permeability and solubility are as follows: CO2 > O2 > N2. This order is consistent with gas condensability, such as in the gas critical temperature: 304 K (CO2), 155 K (O2), and 126 K (N2) (Poling et al. 2000). Meanwhile, the order of the gas diffusivity is as follows: O2 > N2 > CO2. This order is also

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Fig. 4 Gas diffusivity for (a) CO2, (b) N2, and (c) O2 in crystalline and liquid crystalline polymers as a function of crystallinity, Xc. (Reprinted with permission from J Membr Sci., 365, 40. (Copyright (2010) Elsevier)

Polymer PE PTFEP PP PMP PLA PET PS PA PEN HBA/ HNA (LCP) PPTA (LCP)

O2 2.00  10
9 5.98  10
9 4.08  10
10 3.52  10
9 8.26  10
11 4.80  10
12 2.48  10
10 4.57  10
10 2.64  10
12 5.62  10
14

5.61  10
14

Pa CO2 – 5.59  10
8 – 1.18  10
8 2.59  10
10 7.00  10
9 1.50  10
9 1.15  10
9 – 8.01  10
14

2.92  10
12

Crystallinity Range (%) 22–78 31–60 38–62 22–66 4–40 1–51 2–37 24–50 1–39 13–26

22–45

5.00  10
14

N2 7.60  10
10 1.97  10
9 – 9.66  10
10 1.83  10
11 – – 1.38  10
10 – 3.95  10
15 –

Sa CO2 0.550 – – 0.890 2.31 – 2.51 – – 0.072

0.0464

O2 0.150 – – – 0.070 0.095 0.178 – 0.129 0.0060

–

N2 0.100 – – – 0.036 – – – – 0.0012

–

Da CO2 6.76  10
7 – – 2.97  10
6 5.22  10
9 – 4.24  10
8 – – 8.67  10
11

1.02  10
10

O2 1.80  10
6 – – – 4.72  10
8 4.00  10
9 1.04  10
7 – 1.56  10
9 7.09  10
10

Table 2 Purely amorphous transport parameters Pa, Sa, and Da in crystalline and liquid crystalline polymers determined by extrapolation

–

N2 5.05  10
7 – – – 1.37  10
8 – – – – 1.15  10
10

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consistent with the sizes of the gas molecules, such as in the gas critical volume: 73.4 cm3 mol
1 (O2), 90.1 cm3 mol
1 (N2), and 94.1 cm3 mol
1 (CO2) (Poling et al. 2000). Although the gas permeability and solubility in liquid crystalline polymers are lower than those of crystalline polymers, only the gas diffusivity in liquid crystalline polymers is comparable with or lower than that of crystalline polymers like PET and PEN. The ratio of purely amorphous solubility in crystalline and liquid crystalline polymers, Sa, determined via experimentation from Eq. 2 and extrapolation from Fig. 3 is summarized in Table 3. The values of Sa (experimentation) correspond well with Sa (extrapolation) in each gas because there is only a deviation of 20% between the data sets. In contrast, N2 solubility and liquid crystalline polymers show larger deviations compared with other gases and crystalline polymers based on the limited experimental data in this analysis. Therefore, the purely amorphous solubility, Sa, in crystalline and liquid crystalline polymers can be estimated and expressed using the crystallinity and the solubility values in crystalline and liquid crystalline polymers. In addition, the solubility, Sa, can also be estimated from the relationship between the solubility and the crystallinity in the same family of crystalline and liquid crystalline polymers (Fig. 3). The relationship between γ estimated from Eq. 8 and crystallinity is shown in Fig. 5. In this figure, the γ values estimated using Sa (experimentation) or Sa (extrapolation) were plotted. Interestingly, the value of γ is over 1 at crystallinities below 40% for CO2, 55% for O2, and 30% for N2. This result suggests that the diffusivity in crystalline and liquid crystalline polymers can be higher than the purely amorphous diffusivity, Da according to Eq. 4. In particular, the γ of almost all crystalline polymers for CO2 are higher than 1 at the crystallinity range of 0–40%. In general, the gas diffusivity in crystalline and liquid crystalline polymers decreases with increasing crystallinity. However, these results suggest that gas diffusivity at regions of lower crystallinities seems to increase compared with gas Table 3 Ratio of purely amorphous solubility in crystalline and liquid crystalline polymers, Sa, determined by experimentation and extrapolation Polymer PE PTFEP PP PMP PLA PET PS PA PEN HBA/HNA (LCP) PPTA (LCP)

Crystallinity range (%) 22–78 31–60 38–62 22–66 4–40 1–51 2–37 24–50 1–39 13–26 22–45

Experimentation/extrapolation CO2 O2 0.929 0.907 – – – – 1.20 – 1.03 1.20 – 1.12 1.08 0.944 – – – 1.20 1.64 2.07 – 0.506

N2 0.610 – – – 1.25 – – – – – –

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Gas Permeation and Barrier Properties of Liquid Crystalline Polymers

Fig. 5 Diffusivity parameter for (a) CO2, (b) N2, and (c) O2 as a function of crystallinity, Xc. The represented symbols are: crystalline polymers (experimentation) (●), crystalline polymers (extrapolation) (○), liquid crystalline polymers (experimentation) (~), and liquid crystalline polymers (extrapolation) (Δ). (Reprinted with permission from J Membr Sci., 365, 40. (Copyright (2010) Elsevier)

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Fig. 6 Relationship between diffusivity parameter estimated from extrapolation in Fig. 3 and calculated from Eq. 2. The represented symbols are: crystalline polymers for CO2 (experimentation) (●), CO2 (extrapolation) (○), O2 (experimentation) (~), O2 (extrapolation) (Δ), N2 (experimentation) (■), and N2 (extrapolation) (□). (Reprinted with permission from J Membr Sci., 365, 40. (Copyright (2010) Elsevier)

diffusivity in amorphous polymers in this analysis. Meanwhile, the γ of almost all crystalline and liquid crystalline polymers is lower than 1 at crystallinities over 40% for CO2, 55% for O2, and 30% for N2. Based on these results, the crystalline phase in crystalline and liquid crystalline polymers influence gas diffusivity through the following: (1) the crystalline phase can increase gas diffusivity at lower crystallinity ranges, and (2) crystallinity can restrict gas diffusivity at regions of higher crystallinity. Figure 6 shows the comparison between the γ calculated using Sa (experimentation) and Sa (extrapolation), respectively. No significant deviation is observed between Sa (experimentation) and Sa (extrapolation). Therefore, a good agreement between both γ calculated with Sa (experimentation) and Sa (extrapolation) is observed in each gas.

Relationship of Gas Permeability, Solubility, and Diffusivity Figure 7 presents the relationship between gas permeability and the diffusivity in crystalline polymers; the amorphous data are included for comparison (Kanehashi and Nagai 2005). Interestingly, the diffusivity in other crystalline with the exception of PET and PEN and liquid crystalline polymers is higher than that in glassy amorphous polymers at the same gas permeability. The least square fit lines are shown in Fig. 7. The r2 values for CO2 in crystalline, liquid crystalline, and glassy amorphous polymers are 0.559, 0.947, and 0.999, respectively. Similarly, the corresponding r2 values for O2 are 0.764, 0.873, and 0.961, whereas those for N2 are 0.351, 0.942, and 0.902. Glassy amorphous polymers present better correlations compared with the crystalline and liquid crystalline polymers. By contrast, a lower r2

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Gas Permeation and Barrier Properties of Liquid Crystalline Polymers

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Fig. 7 Gas permeability for (a) CO2, (b) N2, and (c) O2 as a function of gas diffusivity. The represented symbols are: crystalline polymers (●), liquid crystalline polymers (~), crystalline PET (■), crystalline PEN (♦), crystalline PS (▼), and conventional amorphous glassy polymers (○). (Reprinted with permission from J Membr Sci., 365, 40. (Copyright (2010) Elsevier)

is obtained in crystalline polymers regardless of the gases. Notably, better correlations are observed in the relationship between gas permeability and diffusivity than the gas permeability and solubility. According to these results, the gas permeability

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Fig. 8 Relationship between gas diffusivity and solubility in the amorphous phase compared to crystalline and liquid crystalline polymers for (a) CO2, (b) N2, and (c) O2. The represented symbols are: crystalline polymers (●), liquid crystalline polymers (~), and conventional amorphous glassy polymers (○). (Reprinted with permission from J Membr Sci., 365, 40. (Copyright (2010) Elsevier)

in crystalline and liquid crystalline polymers seems to depend more on the gas diffusivity rather than the solubility. Furthermore, the crystalline phase effectively increases gas diffusivity in crystalline and liquid crystalline polymers. This result

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Gas Permeation and Barrier Properties of Liquid Crystalline Polymers

543

suggests that the crystalline phase can be attributed to the increase in gas diffusivity in crystalline and liquid crystalline polymers. Figure 8 presents the relationship between gas diffusivity and solubility in crystalline and liquid crystalline polymers, with the amorphous data for comparison (Kanehashi and Nagai 2005). Gas permeability is a product of the solubility and diffusivity, as shown in Eq. 1. Compared with solubility, the gas diffusivities in crystalline, liquid crystalline, and glassy amorphous polymers are widely distributed. For example, O2 diffusivity varies from 3.90  10
9 to 6.70  10
7 for glassy amorphous polymers, 1.10  10
9 to 1.37  10
6 for crystalline polymers, and 6.20  10
11 to 7.65  10
7 cm2 s
1 for liquid crystalline polymers. Oxygen solubility varies from 0.15 to 1.44 for amorphous polymers, 0.00680 to 0.190 for crystalline polymers, and 0.00480 to 0.177 cm3 (STP) (cm3 atm)
1 for liquid crystalline polymers. These results suggest that the permeability of non-condensable gases, such as O2 and N2, in crystalline, liquid crystalline, and glassy amorphous polymers depend more on diffusivity rather than on solubility. Gas solubility in crystalline and liquid crystalline polymers shows lower values, and diffusivity shows higher values than that in glassy amorphous polymers at the same gas permeability (Fig. 8). Thus, the gas transport behavior in crystalline and liquid crystalline polymers can be different from glassy amorphous polymers. This difference is attributed to the presence of crystal structures assuming they follow similar solution/diffusion mechanisms. Therefore, gas transport properties in crystalline and liquid crystalline polymers could be strongly affected by the crystal structure parameters, such as crystal size and distribution.

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Contents Definition . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . Introduction . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . Thermotropic Liquid Crystalline Polymers and In Situ Composites . . . . . . . . . . . . . . . . . . . . . . . . . . . Thermotropic Liquid Crystalline Polymers and Molecular Design . . . . . . . . . . . . . . . . . . . . . . . . In Situ Reinforcement . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . Thermotropic Liquid Crystalline Polymers with High Flame Retardance . . . . . . . . . . . . . . . . . . . . . Side-Group Phosphorus-Containing TLCP . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . Main-Chain Phosphorus-Containing TLCP . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . Conclusions and Challenges . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . References . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . .

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Abstract

The flame retardation for polymer materials can be easily achieved by blending small-molecular flame retardants. However, traditional small molecule flame retardants exhibit potential drawbacks during application, including migration and blooming of the additives; deterioration of the polymer performance; and potential persistence, bioaccumulation, and toxicity (PBT); etc. High molecular weight polymers have been found to be less accessible by living organisms, thus have an automatically lower PBT profile than small molecules. As a typical kind of highly flame-retardant liquid crystalline polymers (LCP),

L. Chen · Y.-Z. Wang (*) The Collaborative Innovation Center for Eco-Friendly and Fire-Safety Polymeric Materials, National Engineering Laboratory of Eco-Friendly Polymeric Materials (Sichuan), State Key Laboratory of Polymer Materials Engineering, College of Chemistry, Chengdu, China e-mail: [email protected]; [email protected]; [email protected] © Springer Nature Switzerland AG 2020 L. Zhu, C. Y. Li (eds.), Liquid Crystalline Polymers, Polymers and Polymeric Composites: A Reference Series, https://doi.org/10.1007/978-3-030-43350-5_55

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phosphorus-containing LCPs have been proved to be a class of efficient high molecular weight flame retardants, which can overcome the aforementioned drawbacks, and have potential industrial applications to replace some existing small molecular flame retardants. The recent relevant developments of phosphorus-containing LCPs with high flame retardance and the corresponding in situ flame-retardant composites are reviewed in this chapter. Keywords

Liquid crystalline polymer · Flame retardance · Reinforcement · Composite

Definition Highly flame-retardant liquid crystalline polymers (LCPs) are those polymers that are composed of both mesogenic moiety and flame-retardant moiety in those polymer chains, resulting in extremely high flame retardance. As a typical kind of highly flameretardant LCP, phosphorus-containing LCPs have been proved to be a class of efficient polymeric flame retardants, which can overcome the practical drawbacks of smallmolecular flame retardants, such as migration and blooming problems; deterioration of the inherent properties; potential persistence, bioaccumulation, and toxicity; etc. and have potential industrial applications as the inherent flame-retardant materials.

Introduction Thanks to the remarkable combination of mechanical and physical properties, solvent resistance, low weight, and ease of processing, polymer materials have long been used in the daily lives. Due to the inflammability of the most polymer materials, unfortunately, losses of life and possessions caused by the fire associated with the use of these polymer materials have aroused much concern among consumers, manufacturers, and official regulatory bodies (Lyon 1994; Stevens and Mann 1999; Irvine et al. 2000; Chen and Wang 2010). Therefore, the use of flame retardants to reduce combustibility and suppress the smoke or toxic gases of the polymers after ignition and combustion becomes a real imminence to explore flameretardant materials to reduce or even to avoid the fire threats. Generally, two methods have been established to achieve flame retardance for general polymers. One is chemically incorporating reactive flame retardant into polymer chains via copolymerization, branching, or grafting, while the other is physically introducing additive-type flame retardants into the matrices via blending, coating, surface finishing, dyeing, etc. (Lu and Hamerton 2002; Bourbigot and Duquesne 2007; Laoutid et al. 2009). Concerning the processing simplicity and comparatively low cost, adding flame-retardant additives become the simplest and most useful and attractive approach to achieve flame retardance. However, normally the small-molecule organic and/or inorganic flame retardants exhibit

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inferior thermal stability to the polymer matrix during compounding; on the other hand, they are easy to migrate and leach out gradually due to the relatively low surface energy and unsatisfied compatibility with the matrix, which would deteriorate the mechanical properties of the matrices simultaneously. On the other hand, some small-molecule flame retardants can be persistent, bioaccumulative, and toxic (PBT), which has led to calls for their deselection from use. The PBT effects, combined with extensive regulatory schemes in the EU to address electronic wastes and register chemicals, lead to the regulatory banning of some brominated flame retardants, particularly small molecular ones. High molecular weight polymers have been found to be less accessible by living organisms and so have an automatically lower PBT profile when compared to small molecules (Lyon 1994; Stevens and Mann 1999; Irvine et al. 2000; Chen and Wang 2010). Therefore, both scientists and manufacturers have recognized use of polymeric flame retardants instead of small-molecule flame retardants. Phosphorus-containing thermotropic liquid crystalline polymers (P-TLCP), which integrate the advantages of both TLCP and phosphorus-containing materials, have been found to be a class of efficient polymeric flame retardants, which can overcome the aforementioned drawbacks of small-molecular flame retardants (Chen et al. 2014). In this chapter, different TLCPs with high flame retardance, particularly phosphorus-containing TLCPs, either in the side group or in the main chain, are comprehensively reviewed, and the corresponding in situ composite with different thermoplastic matrices are summarized.

Thermotropic Liquid Crystalline Polymers and In Situ Composites Thermotropic Liquid Crystalline Polymers and Molecular Design Thermotropic liquid crystalline polymers (TLCP), of which the liquid crystalline behavior occurs by heating the polymer above its glass transition temperature (Tg) or melting point (Tm), are best known for their good thermal stability (Zhu et al. 2007; Xing et al. 2008), outstanding chemical resistance (Shiota and Ober 1997; Luzny et al. 1999), high stiffness and strength (Ortiz et al. 1998a, b), as well as low linear viscosity in the liquid crystalline state (Heino et al. 1994; Bualek-Limcharoen et al. 2001), which make them become attractive high performance engineering materials for many applications (Hyun et al. 1992; Han and Bhowmik 1997). The incorporation of rigid and extended structure of mesogenic units to the mainchain polymer gives rise to an increase in melting temperature, high modulus, and high strength. These materials also possess higher Tg and mesophase-isotropic transition temperature (Ti). However, TLCPs with high transition temperature are too viscous to flow at the temperature below their Ti that they exhibit poor melt processability. The high transition temperature behavior of the extended structures is related to low transition entropy. Considerable efforts have been devoted to reduce the transition temperatures of LCPs in order to reach more practical conditions for industrial processing, such as introducing flexible spacers (Ignatious et al. 1995; Jeong et al. 2006), using isomerious monomers as mesogenic groups (Percec and

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Fig. 1 Schematic diagram showing the molecular design routes for reducing the transition temperature of TLCPs

Yourd 1989; Percec and Tsuda 1990; Percec and Kawasumi 1991), introducing nonlinear or kinked monomers (Lin and Hong 2000; Chang et al. 2002), random copolymerization of monomers, as well as introducing bulky substitute onto mesogenic groups (Percec et al. 1984, 1987; Cai and Samulski 1994; Desrosiers et al. 1996; Han et al. 1997), hyperbranched topologies, and dendritic architectures (Percec and Kawasumi 1992; Percec et al. 1994, 1995). Figure 1 gives the schematic diagram showing the most popular molecular design routes for the synthesis of LCPs containing rigid mesogenic units (Donald and Windle 1992).

In Situ Reinforcement TLCPs can be processed and molded to structural articles by means of the extrusion, injection molding, and melt spinning above their isotropic temperatures. Because of the high cost of the mesogenic monomer synthesis and polymerization, unfortunately, TLCPs are far more expensive than the general purpose engineering plastics. As a consequence, it is more cost effective to create polymer composites with superior mechanical performance using TLCP as a minor blending component. Generally, the presence of a high modulus and high strength TLCP as a dispersed phase in an engineering plastic can act as a “processing aid” to reduce the viscosity of the matrix (Supattra et al. 2009), and mesogenic units of TLCP promote a high degree of molecular alignment in the isotropic state. At the same time, their long relaxation times allow the orientation of the chains to be easily frozen in the solidstate, giving rise to the “in situ” formation of microfibrillar structure under certain processing conditions, and microfibrils of the TLCP significantly affect the mechanical properties of the blended materials, for instance, reinforcement (Kiss 1987). The Latin phrase “in situ” stands for “in position” or “on site” literally. In this regard, in situ reinforced composites are much similar to the short-fiber-reinforced composites (Tiong 2003), such as glass fiber, cabon fiber, and mineral whiskers.

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Fig. 2 Development of TLCP phase from disperse droplets into micro-fibrils

Figure 2 illustrates the morphology development of TLCP phase in the thermoplastic matrix under proper processing conditions; among them, shearing and friction between TLCP (disperse phase) and matrix (continuous phase) are most considered. At the beginning, TLCP phase is deformed as ellipsoids or spherical droplets firstly, then the droplets begin to split into smaller ones; finally micro-fibrils with high aspect ratio are obtained. Generally, a rigid polymer shows a positive enthalpy value as it is blended with a flexible-chain polymer, and the small increase in entropy due to the blending in these two polymers is not able to compensate for the enthalpy effect. Unfortunately, molecular chains of main-chain TLCP show a very stiff and rigid-rod nature. In this regard, the free energy of blending TLCP with thermoplastic matrix is therefore positive. That is to say, the compatibility between TLCPs and thermoplastic matrices is not favorable in thermodynamics (Tiong 2003). Basic understanding of several aspects involved in the processing is crucial to develop the in situ reinforced composites with expected mechanical performances. These aspects include rheology, compatibility, crystallization, and the processing-structure-property relationship of the TLCP/thermoplastic blends. Practically, the former two are mostly considered.

Rheology TLCPs display apparent viscosities, which can be one or two orders of magnitude lower than those of conventional thermoplastics. The melt of TLCPs contains much more domains and are more viscous than small molecule nematics (Fig. 3) (Cogswell 1985). At the very beginning of shearing, the polydomain morphology of TLCPs has a high resistance to flow. To overcome the domain structure, a certain stress level must be exceeded initially. Once the material starts to flow, progressive shear thinning prevails at low stress region. After this, the material flows with a higher viscosity because the domains are broken down into smaller sizes with a larger surface area. At higher stress levels, shear thinning predominates again due to the formation of monodomain or homogeneous continuous phase structure (Cogswell 1985). When TLCP and thermoplastic matrices are compounded, anisotropic mesogenic

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Fig. 3 Relationship between the morphology and rheology of TLCP (Cogswell 1985)

moieties of TLCP become oriented along the flow field within isotropic thermoplastic polymer liquid. The flow-induced orientation results in shear thinning viscosity at low shear rates and low melt viscosities. Thus, the fibrillation, morphology, and distribution of TLCP dispersed phase in the matrix is greatly affected by the processing conditions. Moreover, other factors such as viscosity ratio of the components, TLCP content, interfacial adhesion between the components, and the rheological characteristics of the matrix also play a crucial role in the TLCP fibrillation (Tiong 2003). Generally, the most important rheological parameter that regulates the morphology of TLCP in the in situ composites is the viscosity ratio of TLCP to the matrix, which is defined as λ ¼ ηd =ηm where ηd and ηm are the viscosities of TLCP and the matrix, respectively. Generally, a viscosity ratio  1 is a necessary condition for the fibrillation of TLCP.

Compatibility Since the mixing of a rigid TLCP with a flexible-chain polymer is not favorable in thermodynamics, phase separation of the TLCP blend occurs during processing. Therefore, the reinforcing effect of TLCP is lower than that expected from the rule of mixtures. For effective stress transferring from the polymer matrix to TLCP fibrils, a strong interfacial between reinforcing fibrils and the matrix is needed. Compatibilization can also be promoted by molecular interchange reactions between components, such as transesterification between two thermopolyesters. Chen et al. prepared in situ composites with the polycarbonate/acrylonitrile-butadiene-styrene (PC/ABS, 4:1 in weight ratio) and a phosphorus-containing TLCP named PHBDET (Chen et al. 2009a, b). The results suggested compatibility between PHBDET and PC could be controlled by different degree of transesterification

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(where ABS could simply be considered as an inert component), and the intramolecular reaction could be affected by the processing conditions, such as blending temperature and duration. However, tensile properties of the in situ composites are not linearly dependent on the degree of transesterification. The improved compatibility was not always favorable for the microfibrillation of PHBDET in PC-ABS (Chen et al. 2009b). Additionally, the authors proved that a certain extent of transesterification showed a positive influence on the tensile properties of the composites by enhancing the interfacial adhesion between PC and PHDDT phases. By applying a proper processing condition, the composites with expected in situ reinforcement could be achieved.

Thermotropic Liquid Crystalline Polymers with High Flame Retardance Although the majority of the commercial TLCPs are highly resistant to fire due to their high aromatic constitution and charring tendency thereof, which allows them as the inherent flame-retardant engineering materials in many fields, as flame retardants, these TLCPs should be further modified. Introducing the flame-retardant elements (phosphorus for instance) into the molecules of TLCP before blending the TLCP into the required polymer matrix to make up the in situ composites is supposed to be a good way to resolve the contradiction between flame retardance and mechanical properties of traditional flame-retardant materials. Also, these TLCPs with flame-retardant elements considered as polymeric flame retardants exhibit much lower PBT profile when compared to small molecules. The key factor to approaching an expected in situ composite with both flame retardance and reinforcement is to design a TLCP with suitable transition temperature (to meet the processing temperature of the matrices), appropriate compatibility (between TLCP and matrix), and adequate flame-retardant monomer content (to achieve a desirable flame retardance). For TLCP with high flame retardance, phosphorus-containing compounds containing reactive functional groups that can be copolymerized either in the side group or in the main chain are widely used to enhance the flame retardance of TLCPs (Chen et al. 2014).

Side-Group Phosphorus-Containing TLCP Among the phosphorus-containing moieties, DOPO (9,10-dihydro-9-oxa-10phosphaphenanthrene 10-oxide, Scheme 1) is the most commonly used one both as precursor for a flame-retardant additive and as an integrative flame-retardant monomer. DOPO is a phosphaphenanthrene homologue, derived from o-phenylphenol and phosphorus trichloride. The first synthesis of DOPO was reported by Sanko Chemical (Japan) in 1972 (Saito 1972) and originally this compound was designed and utilized for Toyobo’s flame-retardant polyester fibers and textiles. The reactive P–H bond of DOPO can easily react with activated unsaturated bonds, such as acrylates (C=C),

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Scheme 1 Chemical structure (a) and 3D model (b) of DOPO

Scheme 2 Chemical structure of PLCPAr, where X and 1-X denote the overall composition, not the block length (Wang et al. 2002; Chen et al. 2002)

benzoquinone (aromatic C=C), aldehydes (C=O), etc. or epoxides, which is becoming most of the researchers and manufacturers’ first choice (Chang and Chang 1999). DOPO can also be used for designing side-group phosphorus-containing monomers to prepare flame-retardant TLCP. Wang and his co-workers synthesized a series of DOPO-based wholly aromatic thermotropic liquid crystalline copolyesters (PLCPAr, Scheme 2) with p-acetoxybenzoic acid ( p-ABA) as mesogenic unit, DOPO-substituent hydroquinone (DOPO-HQ) as flame-retardant moiety, and terephthalic acid (TPA) and isophthalic acid (IPA) as linking spacers by transesterification polycondensation (Wang et al. 2002; Chen et al. 2002). Because of the random copolymerization and the bulky pendent DOPO groups in the polymer chains, all copolyesters exhibited nematic liquid crystalline behaviors, suggesting that only directional order existed. Thermogravimetric analysis (TGA) showed that all the copolyesters exhibited excellent thermal stability initiated at 430  C and the decomposition residue at 640  C in nitrogen were all above 40 wt%. It was confirmed that the incorporation of DOPO groups led to good flame retardance. The first published article focusing on in situ composites with high flame retardance TLCP were by Wang et al. (2003). The high flame retardance TLCP was PLCPAr3:1 derived from acetylated DOPO-HQ, p-acetoxybenzoic acid ( p-AHB), and terephthalic acid (TPA) in the mole ratio 1: 3: 1. The spinning PET/PLCPAr3:1 composites exhibited very interesting flame-retardant results: only loading 2 wt% of PLCPAr3:1 could enhance the LOI value up to 26.4 from 21.3 vol% of neat PET; and the LOI value of the composites further reached 32.4 vol% while loading of 15 wt% of PLCPAr3:1. Meanwhile, no dripping was found while burning the samples. In addition to good flame-retardant performance, the mechanical property of the PET composites was also inspiring: tensile strength increased with increasing content of PLCPAr3:1 in the composites, and the elongation at break was similar to that of PET. It is well-established nowadays that the mechanical property of a binary or ternary polymer blend is very sensitive to its morphological state. In such system, better deformation and microfibrillation of TLCP phase accord well

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with the better tensile strength and modulus of the composites. This behavior is mainly due to the fibrillation of TLCP, which participate and transform the applied stress due to its inherent strength and stiffness. Also, it is reported that the breakage of materials mostly occurred on the interfaces between the dispersed phase and matrix, and the larger interface area is, the more additional energy is needed (Parameswaran and Shukla 2000). Continuatively, by injection molding, Du and her co-workers prepared of PET/PLCPAr3:1 in situ reinforced composites (Du et al. 2005, 2006). The results suggested in situ reinforcement could also be obtained by injection molding, that tensile strength increased by at least 25% when the amount of PLCPAr3:1 reached 8 wt%. However, the reinforcement of the injection molded sample is not as good as that of the spinning one, which is due to the larger drawing ratio of the latter one. Deng and his co-workers investigated the flame-retardant mechanism of PLCPAr3:1 on PET (Deng et al. 2008). The results revealed that the presence of PLCPAr3:1 promoted char formation of PET and enhanced thermal stability of the charring residue, hence delayed the further decomposition of the composite. Elemental distribution in gaseous products, liquid products, and solid residues after pyrolysis showed that phosphorus mainly existed in liquid products and residues during pyrolysis of both PLCPAr3:1 and the relevant composite, rather than in gaseous products, indicating that the main action was in the condensed phase. Consequently, PLCPAr endowed both good flame retardance and better mechanical properties simultaneously to the in situ composite; however, due to the high rigidity of the main chain (wholly aromatic) and the presence of the bulky pendent substituent on the hydroquinone unit, Tg, nematic transition temperature (TLC), and Ti of the copolyesters were very high (Table 1). Therefore, further studies have been investigated on reducing the transition temperatures as well as expanding the application fields of TLCP with high flame retardance.

Introducing Flexible Spacers Introducing flexible spacers, including alkyl, silicone, and ether linkages, is supposed to increase the chain flexibility/mobility and thus to decrease the transition temperature of such polymers. By introducing ethylene glycol-containing flexible spacers into the mesogenic chains of PLCPAr, Zhao and his co-workers synthesized a phosphorus-containing TLCP named P-TLCP-FS (Scheme 3), where TLC decreased from 290  C of the aforementioned PLCPAr to 205  C (Zhao et al. 2008). Moreover, P-TLCP-FS exhibited low and wide mesophase temperature, ranging from 185 to 330  C (Table 2), which could match the processing temperatures of commonly used engineering polymers. Also due to the phosphorus-containing groups, high flame retardance with a limiting oxygen index (LOI) value of 70 vol% and an Underwriters Laboratories 94 (UL-94) V-0 rating could still be maintained, suggesting the potential application for flame-retardant in situ composites. Vlad-Bubulac et al. investigated a series of phosphorus-containing liquid crystalline copolyesters based on terephthaloyl-bis-(4-oxybenzoylchloride) (TOBC), where two preformed ester groups were incorporated as mesogenic unit, and DOPO-substituent 1,4-naphthoquinone (DOPO-NQ) was utilized as phosphorus-

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Table 1 Structure composition, thermal transition temperature, liquid crystalline phase, and decomposition residue of PLCPAr where the molar ratio between TPA and IPA is 1:1 (Wang et al. 2002; Chen et al. 2002)

X 0.80 0.75 0.67 0.60 0.50 0.33

Thermal transition from DSCa Tg ( C) TLC ( C) 184 – 183 290 185 287 188 – 192 – 192 –

Thermal transition from POMb K ! LC 290 288 285 284 279 271

LC phase Nematic Nematic Nematic Nematic Nematic Nematic

Residue (wt%)c 50 41 52 47 49 49

Tg glass transition temperature, TLC liquid crystalline temperature, K solid phase, LC liquid crystalline a Peak temperatures from DSC were taken as the phase transition temperature; () denoted transition not observed b Phase transition temperature taken from POM observation, first heating cycle at a heating rate of 10  C min1 c Residue at 640  C were detected by TGA at a heating rate of 10  C min1 under nitrogen atmosphere

Scheme 3 Chemical structure of P-TLCP-FS, where X, Y, and 1-X-Y denote the overall composition, not the block length (Zhao et al. 2008)

containing moiety. To decrease the transition temperatures, aliphatic diols as flexible spacers were incorporated, as illustrated in Scheme 4 and Table 3. The authors first investigated the LC behaviors of the phosphorus-containing liquid crystalline copolyesters with different length of aliphatic units (Vlad-Bubulac and Hamciuc 2009). Results suggested that, all the copolyesters exhibited good thermal stability with initial decomposition temperature above 375  C, and Tgs in the range of 89–138  C. The degree of crystallinity increases by increasing the number of methylene repeats, and the copolyester which had the lowest isotropic temperature, was the polymer containing the longest flexible 12-methylene spacer linkage. As for the copolyesters containing shorter flexible spacers (say, 2-, 3-, 4- or 6-methylene units), they exhibited the most birefringent LC textures and showed isotropic temperature higher than 280  C. Further, the influence of the content of the aliphatic unit on the phase behavior of the copolyesters was also investigated (where 1,12dodecanediol was utilized as the flexible spacers) (Serbezeanu et al. 2010b). The authors declared that, the copolyesters that contained >30 mol% dodecanediol

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Table 2 Structure composition, thermal transition temperature, liquid crystalline phase, and decomposition residue of P-TLCP-FS (Zhao et al. 2008)

X: Y 0.68: 0.22 0.58: 0.30 0.48: 0.36

Thermal transition from DSCa Ti Tg ( C) ( C) 149 330 156 274 158 270

Thermal transition from POMb K ! LC 205 192 185

LC phase Nematic Nematic Nematic

Residue (wt%)c 35 43 40

Tg glass transition temperature, Ti isotropic temperature, K solid phase, LC liquid crystalline Peak temperatures from DSC were taken as the phase transition temperature b Phase transition temperature taken from POM observation, first heating cycle at a heating rate of 10  C min1 c Residue at 700  C were detected by TGA at a heating rate of 10  C min1 under nitrogen atmosphere a

Scheme 4 Chemical structure of the phosphorus-containing copolyesters with different aliphatic diols, R = (CH2)n–, n = 2, 4, 6, and 12 (Vlad-Bubulac and Hamciuc 2009; Serbezeanu et al. 2010a)

showed smectic phases, while the copolyesters that contained