Thermal Analysis of Textiles and Fibers (The Textile Institute Book Series) [1 ed.] 0081005725, 9780081005729

Table of contents :
Cover
The Textile Institute Book Series
Thermal Analysis of Textiles and Fibers
Copyright
Dedication
Contents
List of contributors
1 Introduction
1.1 Introduction
References
2 Fiber process–structure–property relationships
2.1 Introduction
2.2 Spinning
2.3 Drawing
2.4 Yarn after-processing—heat setting and bulking
References
Further reading
3 Differential scanning calorimetry (DSC) in fiber research
3.1 The basics of differential scanning calorimetry
3.2 Calibration of differential scanning calorimetries
3.2.1 Temperature calibration
3.2.1.1 Temperature calibration on heating
3.2.1.2 Temperature calibration on cooling
3.2.2 Energy (enthalpy) calibration
3.2.3 Specific heat capacity calibration (heat flow calibration)
3.2.3.1 Calibration and determination of heat capacity with conventional differential scanning calorimetry
3.2.3.2 Measurement of heat capacity with modulated temperature differential scanning calorimetry in the quasiisothermal mode
3.3 Differential scanning calorimetry of fibers
References
Further reading
4 Thermogravimetric analysis of fibers
References
Further reading
5 Thermomechanical analysis of fibers
References
Further reading
6 Dynamic mechanical analysis (DMA) in fiber research
References
7 Thermal analysis of natural fibers
7.1 Introduction
7.2 Structure of natural fibers
7.2.1 Silk
7.2.2 Keratins
7.2.3 Celluloses
7.2.4 Collagens
7.2.5 Other natural fibers
7.3 Thermal analysis of natural fibers
7.3.1 Measurement of crystallinity of natural fibers by differential scanning calorimetry
7.3.1.1 Glass transition
7.3.1.2 Blending method
7.3.2 Thermal analysis characterization of the melting of natural fibers
7.3.3 Thermal analysis measurement of the specific heat capacity of natural fibers
7.3.3.1 Theoretical prediction
7.3.3.2 Experimental measurement
7.3.4 Thermal analysis characterization of the crystallization kinetics of natural polymers
7.3.5 Thermal analysis study of the natural polymer–bound water systems
7.4 Thermal analysis study of the natural polymer–metal ions systems
7.5 Conclusion
References
8 Polyester fibers
8.1 Introduction
8.2 Poly(ethylene terephthalate) history
8.3 Poly(ethylene terephthalate) polymerization
8.4 Characterization of poly(ethylene terephthalate) chip
8.5 Poly(ethylene terephthalate) fiber processing
8.6 Physical properties of poly(ethylene terephthalate)
8.7 Other polyesters
8.8 Thermal analysis (TA) of polyester fibers
8.9 Polyester fiber thermal analysis in the 21st century
8.10 Other polyester fibers, polypropylene terephthalate, poly(butylene terephthalate), polyethylene naphthalate
8.11 Conclusion
References
Further reading
9 Thermal properties of aliphatic polyesters
9.1 An introduction to aliphatic polyesters
9.2 Thermal properties of polyglycolides
9.2.1 Homopolymers
9.2.2 Copolymers
9.2.3 Effect of additives on thermal properties
9.3 Thermal properties of polylactic acid/poly(lactide)s
9.3.1 Homopolymers
9.3.2 Copolymers
9.3.3 Effect of additives on thermal properties
9.4 Thermal properties of polycaprolactones
9.4.1 Homopolymers
9.4.2 Copolymers
9.4.3 Effect of additives on thermal properties
9.5 Thermal properties of polyhydroxyalkanoates
9.5.1 Homopolymers
9.5.2 Copolymers
9.5.3 Effect of additives on thermal properties
9.6 Conclusion
References
Further reading
10 Poly(ethylene naphthalate) [poly(ethylene-2,6-naphthalene dicarboxylate)]
References
11 Polyethylene fibers
References
Further reading
12 Polypropylene fibers
12.1 Introduction
12.2 Manufacturing of polypropylene fibers
12.3 Processing of polypropylene
12.4 Properties of polypropylene fibers
12.5 Thermal analysis of polypropylene fibers
12.5.1 Melting
12.5.2 The glass transition
12.5.3 Comparison of results/experimental conditions
12.6 Conclusion
References
Further reading
13 Thermal analysis of aliphatic nylon fibers
13.1 Nylon fiber production and basics
13.2 Thermal analysis basics of nylon and its fibers
13.3 Nylon-6
13.4 Nylon-6,6
13.5 Other aliphatic nylon fibers
References
14 Thermal analysis of poly(aryl ether ketone) fibers
14.1 Introduction
14.2 Poly(aryl ether ketone) synthesis
14.3 Producers
14.4 Thermal analysis basics of poly(ether ether ketone) and poly(ether ketone ketone)
14.5 Poly(ether ether ketone) fiber
14.6 Poly(ether ketone ketone) fiber
14.7 Conclusion
References
15 Surgical sutures
References
Further reading
16 Thermal analysis of poly(vinyl alcohol) fibers
16.1 Introduction
16.2 Manufacturing of poly(vinyl alcohol) fibers
16.2.1 Wet spinning of poly(vinyl alcohol) fibers
16.2.2 Dry spinning of poly(vinyl alcohol) fibers
16.2.3 Melt spinning of poly(vinyl alcohol) fibers
16.2.4 Gel spinning of poly(vinyl alcohol) fibers
16.2.5 Hot drawing and stabilization of poly(vinyl alcohol) fibers
16.3 Thermal analysis of poly(vinyl alcohol) fibers
16.4 Conclusion
References
Further reading
17 Polybenzimidazole fiber
References
Further reading
18 Thermal analysis of acrylic and carbon fibers
18.1 Overview
18.2 Introduction
18.3 Production of polyacrylonitrile fibers
18.4 Thermal analysis of polyacrylonitrile and polyacrylonitrile fibers
18.4.1 Thermal analysis of polyacrylonitrile
18.4.2 Thermal analysis of polyacrylonitrile fibers
18.4.2.1 Differential scanning calorimetry characterization
TGA measurements with various heating rates in air and nitrogen
TMA measurements
DMA characterization
18.5 Oxidative stabilization of polyacrylonitrile fibers
18.5.1 Effect of chemical composition
18.5.2 Effect of processing parameters
18.5.3 Study of the stabilization by Fourier-transform infrared spectroscopy (FT-IR)
References
Further reading
19 Thermal analysis of liquid crystalline polymers
19.1 Introduction
19.2 Chemical structure of liquid crystalline polymers
19.3 Process–structure–property relationships of liquid crystalline polymer fibers
19.4 Thermal analysis of liquid crystalline polymer fibers
References
Further reading
20 Thermal analysis of temperature responsive fibrous materials
20.1 Introduction
20.2 Working principle of phase-change material
20.3 Types of phase-change material
20.3.1 Hydrated inorganic salt
20.3.2 Organic hydrocarbons
20.3.3 Polyethylene glycol
20.4 Phase-change material incorporated fibers and their thermal management performance
20.5 Applications of phase-change material and its incorporated fibers
20.6 Conclusions and future research
References
Further reading
21 Thermal characterization of fire-protective fabrics
21.1 Introduction
21.2 Measurement of fire-protective performance of fabrics
21.2.1 Determination of thermal stability of fibers or fabrics
21.2.1.1 Transitional characteristics test
21.2.1.2 Flammability test
21.2.2 Determination of thermal insulation capacity of fabrics
21.2.2.1 Radiant-heat insulation test
21.2.2.2 Gas flame insulation test
21.2.2.3 Combined radiant-heat and gas flame insulation test
21.2.2.4 Flash fire insulation tests
21.2.2.5 Liquid flame insulation test
21.2.2.6 Hot surface contact insulation test
21.3 Evaluation of fire-protective performance of fabrics
21.3.1 Thermal stability
21.3.1.1 Transitional characteristics
21.3.1.2 Flammability
21.3.2 Thermal insulation capacity
21.3.2.1 Insulation from radiant-heat
21.3.2.2 Insulation from gas flame
21.3.2.3 Insulation from combined radiant-heat and gas flame
21.3.2.4 Insulation from flash fire
21.3.2.5 Insulation from liquid flame
21.3.2.6 Insulation from hot surface contact
21.4 Key issues related to the fire-protective performance of fabrics
21.5 Summary and conclusion
References
Acronyms and abbreviations
Abbreviations
Acronyms
Index
Back Cover

Citation preview

Thermal Analysis of Textiles and Fibers

The Textile Institute Book Series Incorporated by Royal Charter in 1925, The Textile Institute was established as the professional body for the textile industry to provide support to businesses, practitioners, and academics involved with textiles and to provide routes to professional qualifications through which Institute Members can demonstrate their professional competence. The Institute’s aim is to encourage learning, recognize achievement, reward excellence, and disseminate information about the textiles, clothing, and footwear industries and the associated science, design, and technology; it has a global reach with individual and corporate members in over 80 countries. The Textile Institute Book Series supersedes the former “Woodhead Publishing Series in Textiles” and represents a collaboration between The Textile Institute and Elsevier aimed at ensuring that Institute Members and the textile industry continue to have access to high caliber titles on textile science and technology. Books published in The Textile Institute Book Series are offered on the Elsevier website at: store.elsevier.com and are available to Textile Institute Members at a substantial discount. Textile Institute books still in print are also available directly from the Institute’s website at: www.textileinstitute.org To place an order, or if you are interested in writing a book for this series, please contact Matthew Deans, Senior Publisher: [email protected]

Recently Published and Upcoming Titles in The Textile Institute Book Series: New Trends in Natural Dyes for Textiles, Padma Vankar and Dhara Shukla, 978-0-08-102686-1 Smart Textile Coatings and Laminates, William C. Smith, 2nd Edition, 978-0-08-102428-7 Advanced Textiles for Wound Care, 2nd Edition, S. Rajendran, 978-0-08-102192-7 Manikins for Textile Evaluation, Rajkishore Nayak and Rajiv Padhye, 978-0-08-100909-3 Automation in Garment Manufacturing, Rajkishore Nayak and Rajiv Padhye, 978-0-08-101211-6 Sustainable Fibres and Textiles, Subramanian Senthilkannan Muthu, 978-0-08-102041-8 Sustainability in Denim, Subramanian Senthilkannan Muthu, 978-0-08-102043-2 Circular Economy in Textiles and Apparel, Subramanian Senthilkannan Muthu, 978-0-08-102630-4 Nanofinishing of Textile Materials, Majid Montazer and Tina Harifi, 978-0-08-101214-7 Nanotechnology in Textiles, Rajesh Mishra and Jiri Militky, 978-0-08-102609-0 Inorganic and Composite Fibers, Boris Mahltig and Yordan Kyosev, 978-0-08-102228-3 Smart Textiles for In Situ Monitoring of Composites, Vladan Koncar, 978-0-08-102308-2 Handbook of Properties of Textile and Technical Fibres, 2nd Edition, A. R. Bunsell, 978-0-08-101272-7 Silk, 2nd Edition, K. Murugesh Babu, 978-0-08-102540-6

The Textile Institute Book Series

Thermal Analysis of Textiles and Fibers Edited by

Michael Jaffe Joseph D. Menczel

Woodhead Publishing is an imprint of Elsevier The Officers’ Mess Business Centre, Royston Road, Duxford, CB22 4QH, United Kingdom 50 Hampshire Street, 5th Floor, Cambridge, MA 02139, United States The Boulevard, Langford Lane, Kidlington, OX5 1GB, United Kingdom Copyright © 2020 Elsevier Ltd. All rights reserved. No part of this publication may be reproduced or transmitted in any form or by any means, electronic or mechanical, including photocopying, recording, or any information storage and retrieval system, without permission in writing from the publisher. Details on how to seek permission, further information about the Publisher’s permissions policies and our arrangements with organizations such as the Copyright Clearance Center and the Copyright Licensing Agency, can be found at our website: www.elsevier.com/permissions. This book and the individual contributions contained in it are protected under copyright by the Publisher (other than as may be noted herein). Notices Knowledge and best practice in this field are constantly changing. As new research and experience broaden our understanding, changes in research methods, professional practices, or medical treatment may become necessary. Practitioners and researchers must always rely on their own experience and knowledge in evaluating and using any information, methods, compounds, or experiments described herein. In using such information or methods they should be mindful of their own safety and the safety of others, including parties for whom they have a professional responsibility. To the fullest extent of the law, neither the Publisher nor the authors, contributors, or editors, assume any liability for any injury and/or damage to persons or property as a matter of products liability, negligence or otherwise, or from any use or operation of any methods, products, instructions, or ideas contained in the material herein. British Library Cataloguing-in-Publication Data A catalogue record for this book is available from the British Library Library of Congress Cataloging-in-Publication Data A catalog record for this book is available from the Library of Congress ISBN: 978-0-08-100572-9 (print) ISBN: 978-0-08-100581-1 (online) For information on all Woodhead Publishing publications visit our website at https://www.elsevier.com/books-and-journals

Publisher: Matthew Deans Acquisitions Editor: Brian Guerin Editorial Project Manager: Fernanda A. Oliveira Production Project Manager: Maria Bernard Cover Designer: Victoria Pearson Typeset by MPS Limited, Chennai, India

Dedication

The editors dedicate this work to Professor Vincent B.F. Mathot upon his retirement from the Department of Chemistry at Katholieke Universiteit Leuven for his many contributions to polymer crystallization and thermal analysis, particularly fast scanning calorimetry. His impact on our field is recognized and acknowledged.

Contents

List of contributors 1.

2.

3.

4.

5.

xiii

Introduction Michael Jaffe and Joseph D. Menczel 1.1 Introduction References

1

Fiber process structure property relationships Michael Jaffe, Anthony J. East and Xianhong Feng 2.1 Introduction 2.2 Spinning 2.3 Drawing 2.4 Yarn after-processing—heat setting and bulking References Further reading

7

1 5

7 8 13 14 14 15

Differential scanning calorimetry (DSC) in fiber research Joseph D. Menczel and William S. Kohl 3.1 The basics of differential scanning calorimetry 3.2 Calibration of differential scanning calorimetries 3.2.1 Temperature calibration 3.2.2 Energy (enthalpy) calibration 3.2.3 Specific heat capacity calibration (heat flow calibration) 3.3 Differential scanning calorimetry of fibers References Further reading

17

Thermogravimetric analysis of fibers Joseph D. Menczel References Further reading

71

Thermomechanical analysis of fibers Joseph D. Menczel and Michael Jaffe References Further reading

81

17 23 23 30 31 36 67 69

79 79

93 94

viii

6.

7.

8.

Contents

Dynamic mechanical analysis (DMA) in fiber research Joseph D. Menczel References Thermal analysis of natural fibers Ye Xue, Wenbing Hu and Xiao Hu 7.1 Introduction 7.2 Structure of natural fibers 7.2.1 Silk 7.2.2 Keratins 7.2.3 Celluloses 7.2.4 Collagens 7.2.5 Other natural fibers 7.3 Thermal analysis of natural fibers 7.3.1 Measurement of crystallinity of natural fibers by differential scanning calorimetry 7.3.2 Thermal analysis characterization of the melting of natural fibers 7.3.3 Thermal analysis measurement of the specific heat capacity of natural fibers 7.3.4 Thermal analysis characterization of the crystallization kinetics of natural polymers 7.3.5 Thermal analysis study of the natural polymer bound water systems 7.4 Thermal analysis study of the natural polymer metal ions systems 7.5 Conclusion References Polyester fibers Michel Jaffe, Anthony J. Easts and Xianhong Feng 8.1 Introduction 8.2 Poly(ethylene terephthalate) history 8.3 Poly(ethylene terephthalate) polymerization 8.4 Characterization of poly(ethylene terephthalate) chip 8.5 Poly(ethylene terephthalate) fiber processing 8.6 Physical properties of poly(ethylene terephthalate) 8.7 Other polyesters 8.8 Thermal analysis (TA) of polyester fibers 8.9 Polyester fiber thermal analysis in the 21st century 8.10 Other polyester fibers, polypropylene terephthalate, poly(butylene terephthalate), polyethylene naphthalate 8.11 Conclusion References Further reading

95 103 105 105 106 107 108 108 109 109 109 111 115 117 119 122 124 126 127 133 133 135 135 136 137 138 138 139 145 147 147 147 149

Contents

9.

Thermal properties of aliphatic polyesters Mazeyar Parvinzadeh Gashti, Marlon Bustos, Hicham Alayan, Roya Jamarani and Milan Maric 9.1 An introduction to aliphatic polyesters 9.2 Thermal properties of polyglycolides 9.2.1 Homopolymers 9.2.2 Copolymers 9.2.3 Effect of additives on thermal properties 9.3 Thermal properties of polylactic acid/poly(lactide)s 9.3.1 Homopolymers 9.3.2 Copolymers 9.3.3 Effect of additives on thermal properties 9.4 Thermal properties of polycaprolactones 9.4.1 Homopolymers 9.4.2 Copolymers 9.4.3 Effect of additives on thermal properties 9.5 Thermal properties of polyhydroxyalkanoates 9.5.1 Homopolymers 9.5.2 Copolymers 9.5.3 Effect of additives on thermal properties 9.6 Conclusion References Further reading

10. Poly(ethylene naphthalate) [poly(ethylene-2,6-naphthalene dicarboxylate)] Joseph D. Menczel References

ix

151

151 153 153 156 158 158 158 162 164 166 166 170 172 174 174 175 176 183 183 189

191 196

11. Polyethylene fibers Joseph D. Menczel and Tonson Abraham References Further reading

197

12. Polypropylene fibers Alfre´d Menyha´rd, Joseph D. Menczel and Tonson Abraham 12.1 Introduction 12.2 Manufacturing of polypropylene fibers 12.3 Processing of polypropylene 12.4 Properties of polypropylene fibers 12.5 Thermal analysis of polypropylene fibers 12.5.1 Melting 12.5.2 The glass transition 12.5.3 Comparison of results/experimental conditions

205

203 203

205 207 207 208 211 211 215 216

x

Contents

12.6 Conclusion References Further reading

220 220 222

13. Thermal analysis of aliphatic nylon fibers Lawrence Judovits 13.1 Nylon fiber production and basics 13.2 Thermal analysis basics of nylon and its fibers 13.3 Nylon-6 13.4 Nylon-6,6 13.5 Other aliphatic nylon fibers References

223

14. Thermal analysis of poly(aryl ether ketone) fibers Lawrence Judovits 14.1 Introduction 14.2 Poly(aryl ether ketone) synthesis 14.3 Producers 14.4 Thermal analysis basics of poly(ether ether ketone) and poly(ether ketone ketone) 14.5 Poly(ether ether ketone) fiber 14.6 Poly(ether ketone ketone) fiber 14.7 Conclusion References

247

223 225 232 239 242 243

247 248 248 249 251 254 257 257

15. Surgical sutures Joseph D. Menczel References Further reading

259

16. Thermal analysis of poly(vinyl alcohol) fibers Zohar Ophir 16.1 Introduction 16.2 Manufacturing of poly(vinyl alcohol) fibers 16.2.1 Wet spinning of poly(vinyl alcohol) fibers 16.2.2 Dry spinning of poly(vinyl alcohol) fibers 16.2.3 Melt spinning of poly(vinyl alcohol) fibers 16.2.4 Gel spinning of poly(vinyl alcohol) fibers 16.2.5 Hot drawing and stabilization of poly(vinyl alcohol) fibers 16.3 Thermal analysis of poly(vinyl alcohol) fibers 16.4 Conclusion References Further reading

271

269 270

271 272 272 273 273 273 274 275 289 289 290

Contents

xi

17. Polybenzimidazole fiber Joseph D. Menczel References Further reading

291

18. Thermal analysis of acrylic and carbon fibers Shahram Arbab and Joseph D. Menczel 18.1 Overview 18.2 Introduction 18.3 Production of polyacrylonitrile fibers 18.4 Thermal analysis of polyacrylonitrile and polyacrylonitrile fibers 18.4.1 Thermal analysis of polyacrylonitrile 18.4.2 Thermal analysis of polyacrylonitrile fibers 18.5 Oxidative stabilization of polyacrylonitrile fibers 18.5.1 Effect of chemical composition 18.5.2 Effect of processing parameters 18.5.3 Study of the stabilization by Fourier-transform infrared spectroscopy (FT-IR) References Further reading

297

19. Thermal analysis of liquid crystalline polymers Michael Jaffe, Anthony J. East and Xianhong Feng 19.1 Introduction 19.2 Chemical structure of liquid crystalline polymers 19.3 Process structure property relationships of liquid crystalline polymer fibers 19.4 Thermal analysis of liquid crystalline polymer fibers References Further reading

325

20. Thermal analysis of temperature responsive fibrous materials Danmei Sun, Kashif Iqbal and Muhammad Owais Raza Siddiqui 20.1 Introduction 20.2 Working principle of phase-change material 20.3 Types of phase-change material 20.3.1 Hydrated inorganic salt 20.3.2 Organic hydrocarbons 20.3.3 Polyethylene glycol 20.4 Phase-change material incorporated fibers and their thermal management performance 20.5 Applications of phase-change material and its incorporated fibers 20.6 Conclusions and future research References Further reading

335

295 295

297 298 301 301 301 303 312 313 315 317 321 323

325 326 328 329 333 334

335 336 337 338 338 341 342 347 350 350 353

xii

Contents

21. Thermal characterization of fire-protective fabrics Sumit Mandal , Sabyasachi Gaan, Martin Camenzind, Simon Annaheim and Rene´ M. Rossi 21.1 Introduction 21.2 Measurement of fire-protective performance of fabrics 21.2.1 Determination of thermal stability of fibers or fabrics 21.2.2 Determination of thermal insulation capacity of fabrics 21.3 Evaluation of fire-protective performance of fabrics 21.3.1 Thermal stability 21.3.2 Thermal insulation capacity 21.4 Key issues related to the fire-protective performance of fabrics 21.5 Summary and conclusion References

355

Acronyms and abbreviations Index

389 393

355 358 358 361 368 368 374 383 384 384

List of contributors

Tonson Abraham Retired from ExxonMobil Chemical Company, Avon, OH, United States Hicham Alayan Department of Chemical Engineering, McGill University, Montre´al, QC, Canada Simon Annaheim Empa, Swiss Federal Laboratories for Materials Science and Technology, Laboratory for Biomimetic Membranes and Textiles, St. Gallen, Switzerland Shahram Arbab Textile Engineering Department, ATMT Research Institute, Amirkabir University of Technology, Tehran, Iran Marlon Bustos Department of Chemical Engineering, McGill University, Montre´al, QC, Canada Martin Camenzind Empa, Swiss Federal Laboratories for Materials Science and Technology, Laboratory for Biomimetic Membranes and Textiles, St. Gallen, Switzerland Anthony J. Easts Consultant, Madison, NJ, United States Xianhong Feng Beckton Dickinson and Company, Franklin Lakes, NJ, United States Sabyasachi Gaan Empa, Swiss Federal Laboratories for Materials Science and Technology, Laboratory for Advanced Fibers, St. Gallen, Switzerland Mazeyar Parvinzadeh Gashti Department of Chemical Engineering, McGill University, Montre´al, QC, Canada; Research and Development Laboratory, PRE Labs Inc., Kelowna, BC, Canada Wenbing Hu Department of Polymer Science and Engineering, Nanjing University, Nanjing, P.R. China Xiao Hu Department of Physics and Astronomy, Rowan University, Glassboro, NJ, United States; Department of Molecular & Cellular Biosciences, Rowan University, Glassboro, NJ, United States

xiv

List of contributors

Kashif Iqbal Department of Textile processing, National Textile University, Faisalabad, Pakistan Michael Jaffe New Jersey Innovation Institute, University Heights, Newark, NJ, United States Roya Jamarani Department of Chemical Engineering, McGill University, Montre´al, QC, Canada Lawrence Judovits Arkema, King of Prussia, PA, United States William S. Kohl TA Instruments, New Castle, DE, United States; Present address: Testing Machines, Inc., New Castle, DE, United States Sumit Mandal Empa, Swiss Federal Laboratories for Materials Science and Technology, Laboratory for Biomimetic Membranes and Textiles, St. Gallen, Switzerland Milan Maric Department of Chemical Engineering, McGill University, Montre´al, QC, Canada Joseph D. Menczel Thermal Measurements LLC, Fort Worth, TX, United States Alfre´d Menyha´rd Laboratory of Plastics and Rubber Technology, Department of Physical Chemistry and Materials Science, Budapest University of Technology and Economics, Budapest, Hungary Zohar Ophir New Jersey Innovation Institute, Newark, NJ, United States Rene´ M. Rossi Empa, Swiss Federal Laboratories for Materials Science and Technology, Laboratory for Biomimetic Membranes and Textiles, St. Gallen, Switzerland Muhammad Owais Raza Siddiqui Department of Textile Engineering, NED University of Engineering & Technology, Karachi, Pakistan Danmei Sun School of Textiles and Design, Heriot-Watt University, United Kingdom Ye Xue Department of Physics and Astronomy, Rowan University, Glassboro, NJ, United States; Department of Biomedical Engineering, Rowan University, Glassboro, NJ, United States

Introduction

1

Michael Jaffe1 and Joseph D. Menczel2 1 New Jersey Innovation Institute, Newark, NJ, United States, 2Thermal Measurements LLC, Fort Worth, TX, United States

Abstract Fibers are produced globally in quantities exceeding 100 million tons with about 60% being synthetic fibers. If one were to choose a single technique to characterize fibers, thermal analysis (TA) would be the technique of choice. The critical issues, controlling fiber performance and the parameters, are conveniently measured by TA overlap. Fiber performance is controlled by the total fiber process history and the key structural elements, nano through micro, imparted to the fiber by this history—crystal morphology, crystallinity, molecular orientation, and phase connectivity are all parameters available through TA techniques. In synthetic fiber manufacture, polymer melts or solutions are converted to fibers by the process called spinning—the conversion of a fluid polymer pool to a continuous, highly uniform filament or yarn (yarn is a multiplicity of filaments) by forcing the polymer liquid through a plate with holes (spinneret), cooling the filament (heat transfer) to the solid state (melt spinning), or by a combination of solvent removal (mass transfer) and cooling. Then, this filamentary product can be further modified structurally by stretching in the solid state (process called drawing), annealing at elevated temperatures, and other related processing steps to provide a product with the desired end-use characteristics through precise control of molecular orientation, crystal morphology, and connectivity between ordered and disordered phases. TA is the characterization technique best suited to define fiber structure and properties and, with a knowledge of fiber manufacturing parameters, fiber processing. Once the parametric correlations of fiber structural features and TA are established, TA provides a rapid and accurate technique for a detailed description of fiber morphology, molecular orientation, and, in some cases, phase interconnectivity. The purpose of this book is to provide the fiber scientist and thermal analyst, focused on fiber materials, an in-depth understanding of how the careful application of TA techniques to fiber problems allows the accurate measurement of key fiber parameters.

1.1

Introduction

If one were to choose a single technique to characterize fibers, TA would be the technique of choice. The critical issues, controlling fiber performance and the parameters conveniently measured by TA overlap. Fiber performance is controlled by the total fiber process history and the key structural elements, nano through micro, imparted to the fiber by this history—crystal morphology, crystallinity, molecular orientation, and Thermal Analysis of Textiles and Fibers. DOI: https://doi.org/10.1016/B978-0-08-100572-9.00001-X © 2020 Elsevier Ltd. All rights reserved.

2

Thermal Analysis of Textiles and Fibers

Understand the origins of performance Performance/Application

Industry

Marketing

Structure

Unique fingerprint Total process history

Quantify effects of processing

Properties mean uniformity

Identify performance predictive parameters

Figure 1.1 Process Structure Property Performance Tetrahedron. Source: M. Jaffe, original drawing.

phase connectivity are all parameters available through TA techniques. The materials science tetrahedron—see Fig. 1.1—illustrates this key relationship. Quoting ourselves from the chapter “Fibers”—in the Thermal Analysis of Polymeric Materials, ed. E. Turi, 1981, 1997)—“The thermal analysis of fibers is the thermal analysis of oriented, semi-crystalline polymers” (Jaffe, 1981; Jaffe et al., 1997). This is true for all fibers, natural or synthetic, and all products engineered from fibers, including textile and industrial fabrics, nonwovens, and fiber-reinforced composite structures. The key formation variables defining fiber structure and properties are as follows: G

G

Backbone chemistry Process history - Thermal history - Mass transfer - Stress during processing

In synthetic fiber manufacture, polymer melts or solutions are converted to fibers by the process called spinning—the conversion of a fluid polymer pool to a continuous, highly uniform filament or yarn (yarn is a multiplicity of filaments) by forcing the polymer liquid through a plate with holes (spinneret), cooling the filament (heat transfer) to the solid state (melt spinning), or by a combination of solvent removal (mass transfer) and cooling. Solvent-based spinning processes include dry spinning, wet spinning, dry-jet wet spinning, electrospinning, and centrifugal force spinning. In all cases the stress on the fiber during the cooling process plays a highly significant role in forming the fiber morphology at the nano- and microscales—see Fig. 1.2. Then, this filamentary product can be further modified structurally by stretching in the solid state (process called drawing)—see Fig. 1.3—annealing at elevated temperatures and other related processing steps to provide a product with the desired end-use characteristics through precise control of molecular orientation, crystal morphology, and connectivity between ordered and disordered phases. In natural fibers, biological processes replace the engineering steps previously described but biologically accomplish the same end state of producing fibers with

Spinning processes (b) Dry spinning

(a) Melting spinning

Feed

Feed

Polymer Chips

Filtered polymer solution

(c) Wet spinning

Filtered polymer solution

Feed

Melting Metered extrusion Filter and spinneret

Metered extrusion Filter and spinneret

Metered extrusion Filter and spinneret

Solidification by solvent evaporation

Cooling and solidifying

Coagulation or Regeneration bath

Solidification or solvent removal Heated chamber Stretch Batch process

Bobbin or cake

Packaging Feed roll and guide Wet treating Lubrication Lubrication Washing

Feed rolls Yarn driving Yarn driving Bobbin Packaging

Packaging Bobbin Drive Roll

Dryer

Drying Oven

Figure 1.2 Spinning process. Source: From Billmeyer, F.W., 1984. Textbook of Polymer Science, third ed. Wiley, p. 493, Fig. 18.4/a.

Figure 1.3 Drawing process. Source: From Billmeyer, F.W., 1984. Textbook of Polymer Science, third ed. Wiley, p. 494, Fig. 18.5.

4

Thermal Analysis of Textiles and Fibers

defined molecular orientation, morphology, and crystallinity. Detailed descriptions of fiber formation and the various means of converting fibers to fabrics and composites may be found in a rich literature—see, for example, the Handbook of Textile Fiber Structure, edited by Eichhorn, Hearle, Jaffe, and Kikutani (Eichhorn et al., 2009a,b), and others (Lewin, 2006, 1983a,b,c). In recent years, processes to produce nano-fibers (fibers with diameters of ,1 µm, often in the 100 nm range) have become of interest—see treatments of electrospinning (Doshi and Reneker, 1995; Shin et al., 2001), centrifugal spinning (Zhang and Lu, 2014), and multilayer processing (Boland et al., 2001). The driving forces, to produce nano-fibers, are biological interest in producing substrates for cellular growth similar to the extracellular matrix (Mo et al., 2004; Xu et al., 2004) and the ability to produce fiber-engineered structures with increased surface area, comfort, and unique mechanical properties (Yoshimoto et al., 2003; Matthews et al., 2002). TA is the characterization technique best suited to define fiber structure and properties and, with a knowledge of fiber manufacturing parameters, fiber processing. Fiber structure and properties are defined by detailed fiber micro- and nanostructure, this structure, in turn, is defined by the detailed stress, thermal, and mass change history imparted to the fiber during spinning, drawing, and all subsequent processing steps, including storage and end-use conditions (Eichhorn et al., 2009a,b). Crystallinity, crystal morphology, and crystal uniformity may be monitored by differential thermal analysis (DSC), augmented by wide angle X-ray scattering (WAXS), small angle X-ray scattering (SAXS), optical microscopy (OM), scanning electron microscopy (SEM), transmission electron microscopy (TEM), atomic force microscopy (AFM), and other calibrating data. Molecular orientation may be monitored by thermomechanical analysis (TMA) and dynamic mechanical analysis (DMA), while also providing data about fiber dimensional stability and mechanical properties. Thermogravimetric analysis (TGA) provides data about backbone polymer chemical stability, water uptake, residual solvent concentrations, oxidative stability, and other aspects of environmental stability. Once the parametric correlations of fiber structural features and TA are established, TA provides a rapid and accurate technique for a detailed description of fiber morphology, molecular orientation, and, in some cases, phase interconnectivity. Fibers are produced globally in quantities exceeding 100 million tons with about 60% being synthetic fibers (Carmichael, 2015). Fibers are typically described by their chemistry and linear density (weight/unit length) with the most common units being denier (weight in grams/9000 m) or decitex (weight in grams of 10,000 m). Typical filament deniers are in the range of 2 5, that is, 9000 m of a typical textile filament weighs 5 g or less. Typical yarn deniers are in the range of 100 1000, and variations in filament or yarn deniers (a 5 denier filament is about 12 µm in diameter, depending on fiber chemistry) are about 5%. We leave it to the reader to calculate the length in light-years of global fiber production. Over the decades, fiber producers have developed quality control (QC) techniques to determine yarn performance characteristics, for example, yarn tensile properties are often measured by monitoring the yield parameters of yarn bundles, yarn shrinkage is measured by the shrinkage of hanks of yarn in boiling water, and yarn molecular orientation is often measured by dye

Introduction

5

uptake—a strong function of noncrystalline molecular orientation (Patterson and Ward, 1957; Munden and Palmer, 1950). Wherever possible, fiber and textile manufacture is monitored online with measurements, such as shrinkage tension and online diameter measurement (Uster) (Pegram, 2000). The applications of TA to fiber characterization range from “fingerprinting” to a detailed process structure property investigation of a given fiber formation system. Fingerprinting in this sense means determining that a DSC, TMA, DMA, or TGA responses of a given product are the same, within experimentally defined limits. The mechanistic, morphological, or performance implications of the thermal analysis curves are not evaluated. Coupled with modern, high-speed TA techniques, fingerprinting may be suitable as a QC test. The true strength of TA in fiber science lies in the ability to correlate TA results with morphological detail and/or process variables and end-use properties. TA provides a powerful set of experimental options, for understanding the range of performance, available to a given fiber chemistry and a sophisticated methodology for the diagnosis of fiber manufacturing and performance problems. This book is replete with examples of such studies, as discussed in the chapters on the TA of fibers in earlier Turi books (Jaffe, 1981; Jaffe et al., 1997) and others (Cheng, 2002). The application of TA techniques to fibers requires careful attention to sample geometry and conditioning. The researcher needs to take care not to contribute randomly to the process history of the sample, thus potentially masking the earlier process history. Heating of a fiber may result in increases or changes to the fiber crystallinity or relaxation of fiber molecular orientation. Irreversible changes to fiber structure may be monitored by the application of modulated techniques or through the design of sample holders that keep fibers from relaxing during heating. Fiber slippage in TMA or DMA experiments may also lead to inaccurate and irreproducible results. In DSC experiments, it is sometimes attractive to take the results of a second heating as the fiber response, which is often being more reproducible than the first heating. While this may be appropriate for some experimental outcomes, it should be understood that the first heating of the sample reflects the sample process history, while the second heating is dominated by structural changes that occur during the first heating and may lead to incorrect conclusions—an example of this may be changes noted in Tg of polyamides if the concentration of absorbed water is changed by the experimental history (Batzer and Kreibich, 1981). The purpose of this book is to provide the fiber scientist and thermal analyst, focused on fiber materials, an in-depth understanding of how the careful application of thermal analysis techniques to fiber problems allows the accurate measurement of key fiber parameters.

References Batzer, H., Kreibich, U., 1981. Polym. Bull. 5 (11 12), 585 590. Billmeyer, F.W., 1984. Textbook of Polymer Science, third ed. Wiley. Boland, E.D., Wnek, G.E., Simpson, D.G., Pawlowski, K.J., Bowloin, G.L., 2001. J. Macromol. Sci. 38 (12), 1231 1243.

6

Thermal Analysis of Textiles and Fibers

Carmichael, A.B., 2015. Man-Made Fibers Continue to Grow. Textile World. Cheng, S., 2002. Handbook of Thermal Analysis and Calorimetry: Applications to Polymer and Plastics, vol. 3. Elsevier. Doshi, J., Reneker, D., 1995. J. Electrostat. 35 (2 3), 151 160. Eichhorn, S., Hearle, W.S., Jaffe, M., Kikutani, T., 2009a. Handbook of Textile Fiber Structure, vol. 1. Woodhead Publishing, Oxford, Cambridge, New Delhi. Eichhorn, S., Hearle, W.S., Jaffe, M., Kikutani, T., 2009b. Handbook of Textile Fiber Structure, vol. 2. Woodhead Publishing, Oxford, Cambridge, New Delhi. Jaffe, M., 1981. Fibers. In: Turi, E. (Ed.), Fibers in Thermal Characterization of Polymer Materials. Academic Press, New York. Jaffe, M., Menczel, J.D., Bessey, W.E., 1997. Fibers. In: Turi, E. (Ed.), Thermal Characterization of Polymer Materials, second ed. Academic Press, San Diego, CA. Lewin, M., 2006. Handbook of Fiber Chemistry, third ed. CRC Press. Lewin, M., Sello, S.B., 1983a. Handbook of Fiber Science and Technology, vol. I, Part A. Marcel Dekker, New York. Lewin, M., Sello, S.B., 1983b. Handbook of Fiber Science and Technology, vol. I, Part B. Marcel Dekker, New York. Lewin, M., Sello, S.B., 1983c. Handbook of Fiber Science and Technology, vol. 2. Marcel Dekker, New York. Matthews, J.A., Wnek, G.E., Simpson, D.G., Bowlin, G.L., 2002. Biomacromolecules 3 (2), 232 238. Mo, X.M., Xu, C.Y., Kotaki, M., Ramakrishna, S., 2004. Biomaterials 25 (10), 1883 1890. Munden, A.R., Palmer, H.J., 1950. J. Text. Inst. Proc. 41 (7), 609 634. Patterson, D., Ward, I.M., 1957. Trans. Faraday Soc. 53, 1406 1412. Pegram, J., 2000. Text. Prog. 30 (1 2), 90 97. Shin, Y.M., Hohman, M.M., Brenner, M.P., Rutledge, G.C., 2001. Polymer 42 (25), 09955 09967. Xu, C.Y., Inai, Y., Kotaki, M., Ramakrishna, S., 2004. Tissue Eng. 10 (7 8), 1160 1168. Yoshimoto, H., Shin, Y.M., Terai, H., Vacanti, J.P., 2003. Biomaterials 24 (12), 2077 2082. Zhang, X., Lu, Y., 2014. Polym. Rev. 54 (4), 677 701.

Fiber process structure property relationships

2

Michael Jaffe1, Anthony J. East2 and Xianhong Feng3 1 New Jersey Innovation Institute, Newark, NJ, United States, 2Consultant, Madison, NJ, United States, 3Becton Dickinson and Company, Franklin Lakes, NJ, United States

Abstract The purpose of this discussion of fiber process structure property relationships is to illustrate the power of thermal analysis (TA) in the understanding and characterization of the complex processing history that is reflected in the nano micro macro structures of fibers, as well as the key structural and property parameters that define fiber performance. What is emphasized is the relationship between process conditions, fiber structure formation, and TA techniques. While the examples given relate to specific synthetic fibers with diameters of microns to a few tens of microns, the structure of all fibers (high-aspect-ratio materials of more or less circular cross section) will have similar features, and the relationship of morphology to properties is independent of formation conditions.

2.1

Introduction

The purpose of this discussion of fiber process structure property relationships is to illustrate the power of TA in the understanding and characterization of the complex processing history that is reflected in the nano micro macro structures of fibers, as well as the key structural and property parameters that define fiber performance. In this section, melt spinning will be described in detail as illustrative of all spinning processes. What is emphasized is the relationship between process conditions, fiber structure formation, and TA techniques; for more detail on specific fiber-processing methods or specific fibers see the specific fiber chapters of this book and the rich fiber literature (Eichhorn et al., 2009a,b). Where appropriate, cartoons illustrating structure formation during fiber processing will be employed. The importance of such cartoons is to help one to visualize what happens during the complex and often extremely fast (fiber spinning may reach speeds of .100 mi/h) formation steps that convert bulk polymers to fibers. While the example given relates to synthetic fibers with diameters of microns to a few tens of microns, the structure of all fibers (high-aspect-ratio materials of more or less circular cross section) will have similar features and the relationship of morphology to properties is independent of formation conditions.

Thermal Analysis of Textiles and Fibers. DOI: https://doi.org/10.1016/B978-0-08-100572-9.00002-1 © 2020 Elsevier Ltd. All rights reserved.

8

2.2

Thermal Analysis of Textiles and Fibers

Spinning

The fiber spinning of poly(ethyeneterephthalate) (PET), polypropylene and nylon, and other synthetics and partial synthetics (e.g., cellulose acetate or mercerized cotton) has been extensively treated in the patent literature (e.g., search the fiber patents of DuPont, too numerous to list here), somewhat less so in the open literature, although the book edited by Hearle (Eichhorn et al., 2009a,b) and other reviews (Takajima et al., 1994; Zhang, 2014) are good introductions to the subject. We will concentrate here on how changes in the key process variables of spin-line stress and temperature profile affect assembly at the molecular level through the nano- and microscale (morphology), and in turn, how the morphology affects the resulting performance of the yarn. The relationships described here are equally valid for all semicrystalline polymers. For fiber products to have commercial value, both the average value of key properties and the standard deviation associated with the mean value of key properties must be controlled. In general, variation in properties, hence variation in morphology must be controlled to ,10% for the yarn to be commercially acceptable. Variation means measureable differences between filaments in a yarn or along a given filament. The frequency of variation is also critical, high-frequency changes that may be averaged over a critical use length, are, in general, more acceptable than smaller variation along or between filaments that occurs at the lower frequency. Molten polymer is introduced (or produced) into the manifold of the spinning machine. The manifold may feed as few as one or as many as 200 separate spinnerets and is designed to keep the directed polymer streams as uniform as possible in shear and thermal history. Thermal and shear history leaves an imprint on all TA measurements, but accurate data interpretation requires knowledge of the fiber process history, for example, PET spinning temperature is typically between 280 C and 300 C, but local shear heating may increase this temperature by as much as 10 C 15 C. The molten polymer stream is then fed through metering pumps to the spinning pack (assembly that starts with a series of filters and ends at the spinneret) —see Fig. 2.1. The spinneret consists of between five (hosiery yarn) and several thousand of holes, typically ranging from 180 to 400 µm in diameter. Pack and spinneret design are highly specialized, and the reader is referred to the open and patent literature for the depth of engineering detail available on these subjects. The purpose of the pack and spinneret is to ensure that filtered (clean) polymer is fed to each hole of the spinneret as uniformly as possible. Passage through the spinneret subjects the polymer to a complex rheological environment (see, e.g., the work of Denn, 1996; Kikutani et al., 1996), resulting in local increases of molecular orientation and a distribution of orientation between spinneret wall and centerline. Upon exiting the spinneret the combined effects of surface tension and relaxation of molecular orientation results in die swell (increase of the filament diameter to greater than the spinneret hole diameter), which can be countered by applying a stress leading to a controlled reduction in the fiber diameter (drawdown).

Fiber process structure property relationships

9

The conversion of materials to fibers is the conversion of a liquid to long, thin things

Liquid pool Material synthesis Solidification

Take-up

Modify structure

Modify surface

Figure 2.1 Melt spinning. Source: Jaffe, unpublished figure.

From a molecular point of view the starting polymer melt is an unoriented entangled network (visualize a bowl of cooked noodles), characterized by the polymer molecular weight, molecular-weight distribution, entanglement density, and the average chain length between entanglements. The polymer melt is further complicated by particulate matter that may include catalyst residues, stabilizers, lubricants, and random dust. The orientation imparted during drawdown generally leads to a net increase of molecular orientation of the fiber. Concurrently, entropic relaxation is occurring and orientation is being lost, and a dynamic balance of stretch-induced orientation and molecular relaxation is established. The imparted orientation is not necessarily uniform, and it is likely that a distribution of molecular orientation exists within the spun fiber. The higher the spinning speed or the more viscous the polymer (lower spinning temperatures or higher molecular weight) the less the time available for relaxation and the higher the residual orientation. Thermo mechanical analysis (TMA) shrinkage measurements can give an indication of the amount of amorphous orientation present in fiber or yarn samples—see TMA measurements. Shrinkage force measurements can be correlated with phase interconnectivity (tie molecules) and the orientation of the network—see DMA measurement. The processes that occur in the spin line, between the exit of the polymer from the spinneret to the point of first stress isolation on the first roller (or equivalent fiber capture device) at the base of the spin line, all involve the changing of this fluid molecular network to the solid-state molecular chain topology of the filament. Within a distance of a few meters, and under the influence of the applied force (take-up tension) and quench media, at speeds in excess of 100 mi/h—less than 0.01 second residence time—the fiber is transformed from a fluid molecular network to a highly interconnected semicrystalline morphology, characterized by the amount, size, shape, and net orientation (with respect to the fiber or the long axis)

10

Thermal Analysis of Textiles and Fibers

of crystalline units, and the orientation and spatial distribution of noncrystalline chains. All of these units are interconnected by molecules that traverse more than one local ordered region (tie molecules), the load-bearing elements of the fiber structure. Differential scanning calorimetry (DSC), shrinkage/shrinkage force, and dynamic mechanical analysis (DMA) measurements give insight into the nature and stability of this network. It has long been noted (see Ziabicki, 1976) that the crystallization rate of a polymer may increase by up to six orders of magnitude if the crystallization event occurs when the polymer chains are under an applied stress and have a net orientation (anisotropy) rather than being in a random orientation state (isotropic). This large increase in crystallization rate is accompanied by a change in crystal habit, the shape of the crystalline phase produced transforms, over a narrow stress regime, from a spherulitic (spherically symmetrical) to a columnar habit—see Fig. 2.2—and a significant increase in crystallinity. These changes manifest in the melting behavior of the spun yarn—see, for example, the treatment of polypropylene fiber morphology as a function of spinning stress by Jaffe (1978), or spinning speed by Heuval and Huisman (1981). This morphological transition is surprisingly sharp and occurs at an applied stress of about 0.1 g/d (B7 MPa). Increasing the spin-line stress leads to an increase in the number of rows and a decrease in the diameter of the fibrillar structures. As the fibrils are stable only in the presence of the spinning stress (oriented melt entropy), they may or may not be visible in the final fiber morphology. A useful way of conceptualizing the process is to divide the spin line into the following three regions: G

G

G

Region 1: Increase local and global molecular orientation Region 2: Fibril formation at points of maximum orientation (transient mesogen, mechanical steady state) Region 3: Fibril decoration (folded-chain crystal growth)

Crystal habit, crystallization rate

Spherulite

Row structure “shish-kebab”

Spinning speed, spin-line stress, melt orientation

Figure 2.2 Melt spinning morphology development. Source: Jaffe, unpublished figure.

Fiber process structure property relationships

11

Figure 2.3 Fiber spinning structure formation. Source: Jaffe, unpublished figure.

A cartoon of this model of morphology/molecular chain topology development in melt spinning is shown in Fig. 2.3—see work of Jaffe et al. (Lewin, 2006). The details and experimental validation of these complex processes have remained elusive over the past several decades, and a number of models and suggestions exist in the literature (Meerveld et al., 2008). Until the details of the nucleation process in a spinning fiber are elucidated, however, description of structure formation in the spin line is deduction and conjecture. The value of this model is that it provides a useful framework for visualizing fiber structure formation and can correlate structure with subsequent processing steps and, ultimately, with fiber properties. The model is an effective tool for defining process parameter changes that lead to desired performance changes. In Region 1 the spin-line stress leads to filament drawdown, causing a net increase of molecular orientation of the molten polymer. A consequence of this stress is the disentangling of some of the starting network chains and the increase of local molecular chain orientation in the proximity of remaining entanglements and particles (stress concentration points). As these bundles of locally oriented chains grow in aspect ratio, they may satisfy the conditions for nematic phase formation (Onsager, 1931; Flory, 1942) or perhaps crystallization, leading to a biphasic array comprising fibrillar (mesogenic) structures sitting in a less-oriented amorphous matrix. When the spin-line stress is completely supported by these fibrillar structures, the matrix chains are able to relax, and the conditions for fibril formation no longer exist (Region 2). As one enters the lower temperature ranges, the fibrils provide an effective, high density of nucleation sites for conventional lamellar crystal overgrowth of the fibrils, leading to increases of up to six orders of magnitude in the effective crystallization rate (Region 3). Fibrils may or may not be evident in the final structure, but the high orientation of the wholly semicrystalline structure is always evident. The resulting morphology acts as a template for all further structure formation. Jaffe has shown that these nuclei persist even in the melt

12

Thermal Analysis of Textiles and Fibers

is subjected to temperatures well above the melting point for a long time (Jaffe, 1978). It is also evident that a population of the molecular chains can participate in more than one element of the structure; these tie molecules provide the stresstransfer elements in subsequent fiber deformation (Lewin, 2006). Conceptually, three types of tie molecules are possible in the model—interfibrillar, interlamellar (between lamellae on a given fibril or between lamellae on different fibrils), and between fibril and lamella. It is the tie-molecule distribution, combined with the remaining entanglement distribution that defines the residual draw ratio (see Ward, 1997, 1983) of the fiber structure. The detailed proof of this conceptual model is experimentally difficult, although it is generally supported by existing experimental data and melt-spinning process models. The overall veracity of the model is less important than the utility of the model in predicting process structure property relationships. Important implications of the model are as follows: G

G

G

G

G

The order of molecular chain orientation and crystallization steps in fiber spinning is critical. The formation of an oriented fibrillary structure is the template for all further morphological developments and defines the nucleation density for subsequent crystallization. As chain orientation prior to crystallization is increased, the load-bearing aspects of the crystalline network produced also increases, while the noncrystalline load-bearing elements of the structure decrease. This leads to the decoupling of molecular orientation responsible for increased modulus and strength, from oriented chains responsible for entropic shrinkage, allowing for highmodulus low-shrinkage fiber products. The network defined in spinning remains the template for structure formation in all subsequent processing steps.

The melt spinning of all semicrystalline polymers can be fit into the general framework described earlier. Details of specific melt-spinning processes are well 0.9 IV PET

6

y = 6.6479 + –2.81681log(x) R = 0.99291

DRmax

5 4 3 2 1

0

20

40

60

80

100

120

Spun BI

Figure 2.4 Maximum draw ratio as a function of spun birefringence. Source: Jaffe, unpublished figure.

Fiber process structure property relationships

13

documented in the chapter by Reese et al. (Ward, 1997), East et al. (Lewin, 2006), or the papers of Ward (1997, 1983), Prevorsek (Lewin, 2006), and others (Lewin, 2006). The structural state of the spun yarn, while complex, is often described by a single parameter, the spun-yarn birefringence, an average measure of fiber molecular orientation. It has been shown by Jaffe that the spun-yarn shrinkage is an excellent predictor of the remaining yarn draw ratio as shown in Fig. 2.4 (Lewin, 2006) where DRmax is defined as the highest stable draw ratio available to a given spun yarn. Ward has derived a similar relationship based on rubber elasticity theory (Ward, 1983). Some examples of how this history affects fiber TA response may be found in the work both done and summarized by Jaffe and Menczel (Jaffe, 1981; Jaffe et al., 1997).

2.3

Drawing

Despite the orientation introduced during spinning, additional increases in molecular order are often brought about by a separate drawing process, that is, stretching the fiber in the solid state between rollers rotating at different speeds where the second roller rotates faster than the first and under conditions where the imparted crystalline or molecular orientation does not have time to relax. As-spun fibers can be amorphous or crystalline, depending on the polymer chemistry and the spinning conditions. Most fibers become more crystalline and better oriented when drawn. Faster crystallizing polymers, such as polypropylene (PP), polyethylene (PE), poly (butyleneterephthalate) (PBT), poly(propylenterephthalate) (PPT), or nylon-6,6, always form semicrystalline spun fibers, although they still often need a solid-state drawing stage to induce/complete crystallite orientation. The combination of molecular entanglements and the presence of polymer crystallites lock this orientation into place. This, in turn, affects such parameters as tenacity, modulus, elongation at break, and heat shrinkage. For drawing to be most effective the fiber must be drawn close to its maximum draw ratio. Draw ratio is the ratio of yarn feed velocity to draw-roll haul-off speed: it can vary from about 1.5 up to 6.0 in conventional fibers and as high as 30 1 in superdrawn (Smith and Lemstra, 1980) fibers such as highmodulus polyethylene fibers (Smith et al., 1981). The draw point is the actual place where fiber drawing takes place and must be stabilized. In early processes, this was done by a heated metal snubber pin, usually heated to about Tg 1 10 C, but this tended to produce drawn yarn with an unacceptable degree of heat shrinkage. The latter was minimized by heat setting the fiber by passing it over a long hot plate set well above the effective Tg of the drawn, crystallized yarn. This simple system was adequate when draw speeds were low (,500 m/min), but as draw speeds rose considerably, it was necessary to use separately heated feed rolls and draw rolls to achieve the same effect at much higher speeds. The draw ratio has a major effect on yarn elongation and tenacity. As one would expect, high draw ratios give high tenacity yarns with higher yarn moduli and lower extensions to break; low draw ratios give lower tenacities with higher extension. It was shown by Jaffe

14

Thermal Analysis of Textiles and Fibers

(Lewin, 2006), Ward (1997), and others that a consequence of high-speed spinning is to shift the load supporting of the network chains of the fiber structure from noncrystalline to crystalline regions of the fiber morphology. This limits the draw ratio available to fully orient these fibers, resulting in fibers with nearly equivalent tensile properties but significantly lower shrinkage at elevated temperature. Examination of the DSC, TMA, and DMA of drawn yarns can give insight into the total process of the yarns tested and can allow the separation of spinning effects from drawing.

2.4

Yarn after-processing—heat setting and bulking

Drawn filament yarns can be treated in any number of ways. The yarn may simply be wound onto a yarn package, or twisted on a ring frame or sent for a yarn bulking process such as false-twist bulking. Many apparel yarns need to be textured or “bulked” to give desirable esthetic properties, particularly for cotton blends and women’s wear markets. This may be done during drawing (draw-bulking) or in a separate process. The number of bulking processes is numerous and for more detailed descriptions, see the books edited by Eichhorn et al. (2009a,b). The principle of so-called false-twist bulking is to create minor side-to-side variations in molecular orientation across a given yarn, causing the yarn to bend during controlled thermal shrinkage to create a 3D structure with a bulky feel. After bulking, fibers/yarns undergo many additional processing steps, including dyeing, fabric formation (knitting, weaving), heat setting (annealing under controlled shrinkage conditions), and finally conversion to garments, carpets, industrial fabrics, ropes, and other end uses too numerous to elaborate. Turi (Jaffe, 1981; Jaffe et al., 1997), Valk (Jaffe, 1981; Jaffe et al., 1997), and others investigated the effects of fiber/fabric annealing on the fiber TA. Any process step that subjects the yarn or fiber to heat, solvents or stresses will affect fiber structure and, hence, fiber TA response. Although some reviews of these processes exist (Jaffe, 1981; Jaffe et al., 1997), a detailed description of these processes and their impact on fiber TA is beyond the scope of this book.

References Denn, M., 1996. Industrial & Engineering Chemistry Research 35 (9), 2842 2843. Eichhorn, S., Hearle, W.S., Jaffe, M., Kikutani, T., 2009a. Handbook of Textile Fiber Structure, vol. 1. Woodhead Publishing, Oxford, Cambridge, New Delhi. Eichhorn, S., Hearle, W.S., Jaffe, M., Kikutani, T., 2009b. Handbook of Textile Fiber Structure, vol. 2. Woodhead Publishing, Oxford, Cambridge, New Delhi. Flory, P., 1942. Thermodymamics of high polymer solutions. J. Chem. Phys. 9 (8), 660. Heuval, H.M., Huisman, R., 1981. J. Appl. Polym. Sci. 26 (2), 713 732. Jaffe, M., 1978. In: Shalaby, S.W. (Ed.), Thermal Analysis of As-Spun Polypropylene Fibers. Franklin Institute Press, Philadelphia, PA.

Fiber process structure property relationships

15

Jaffe, M., 1981. Fibers. In: Turi, E. (Ed.), Thermal Characterization of Polymer Materials. Academic Press, New York. Jaffe, M., Menczel, J.D., Bessey, W.E., 1997. Fibers. In: Turi, E. (Ed.), Thermal Characterization of Polymer Materials, second ed. Academic Press, San Diego, CA. Kikutani, T., Radhakrishnan, J., Arikawa, S., Takaku, A., Okui, N., Jin, X., et al., 1996. J. Appl. Polym. 62 (11), 1913 1924. Lewin, M., 2006. Handbook of Fiber Chemistry, third ed CRC Press. Meerveld, J., Hutter, M., Peters, G., 2008. Continuum Model for the Simulation of Fiber Spinning, With Quiescent and Flow-Induced Crystallization 150 (2 3), 177 195. Onsager, L., 1931. Reciprocal relations in irreversible process. Phys. Rev. 38, 2265. Smith, P., Lemstra, P., 1980. Polym. Int. 12 (4), 212 214. Smith, P., Lemstra, P., Booij, H., 1981. J. Polym. Sci., B., Polym. Phys. 19 (5), 877 888. Takajima, T., Kajiwara, K., McIntyre, J.E., 1994. Advanced Fiber Spinning Technology. Woodhead Publishing, Japan. Ward, I.M., 1983. Mechanical Properties of Solid Polymers, second ed. Wiley, NY. Ward, I.M., 1997. Structure and Properties of Oriented Polymer. Springer. Zhang, D., 2014. Advances in Filament Yarn Spinning of Textiles and Polymers. Woodhead Publishing, UK. Ziabicki, A., 1976. Fundamentals of Fibre Formation: The Science of Fibre Spinning and Drawing. Wiley, UK.

Further reading Kilian, J., 1955. Spinneret Assembly, US2936482A, E I du Pont De Nemours and Co. Larrondo, L., John Manley, R., 1981a. J. Polym. Sci., Polym. Phys. 19 (6), 909 920. Larrondo, L., John Manley, R., 1981b. Electrostatic fiber spinning from polymer melts. II. Examination of the flow field in an electrically driven jet. J. Polym. Sci., Polym. Phys. 19 (6), 921 932. Lewin, M., Sello, S., 1983. Handbook of Fiber Science and Technology, vol. 2. Marcel Dekker, New York. Song, W., Kinloch, I., Windle, A., 2003. Science 302 (5649), 1363.

Differential scanning calorimetry (DSC) in fiber research

3

Joseph D. Menczel1 and William S. Kohl2,3 1 Thermal Measurements LLC, Fort Worth, TX, United States, 2TA Instruments, New Castle, DE, United States, 3Present address: Testing Machines, Inc., New Castle, DE, United States

Abstract This chapter describes the fundamentals of differential scanning calorimetry (DSC) and also how DSC can be used for characterization of fibers. The differences between the heat-flux DSC and power-compensation DSC are explained. The various calibrations of DSC (i.e., temperature calibration on heating and cooling, energy and heat capacity calibrations) are also described. The effect of the rigid amorphous and oriented amorphous phases on the glass transition is shown. The change of the phase transitions with increasing orientation is demonstrated. New modulated DSC methods are introduced for describing the various transitions (glass transition, melting, and crystallization). Special emphasis is devoted to superheating. The differences between free-to-shrink and constrained fiber measurements are illustrated for several important synthetic fibers. Examples are given for synthetic and natural fibers.

3.1

The basics of differential scanning calorimetry

By the definition of ASTM (American Society for Testing and Materials, standard E473-04), differential scanning calorimetry (DSC) is a technique in which “the heat flow rate difference into a substance and a reference material is measured as a function of temperature while the substance and reference are subjected to a controlled temperature program.” As Wunderlich (1990) emphasized, no heat flow meter exists for direct measurement of the amount of the heat. Therefore indirect means are needed to measure the heat flowing into or out of the sample. This can be done if we have a reference in addition to the sample, and the temperature difference between them is monitored continuously. In steady state, this temperature difference is proportional to the heat capacity of the sample and the heat flowing into or out of the sample. Two types of DSC techniques have been developed: heat flux and power compensation. G

Heat-flux DSC was evolved directly from differential thermal analysis (DTA). The thermocouples were created by Le Chatelier (1887), and DTA was developed by RobertsAusten (1899) and Kurnakov (1904) in the 18th
19th centuries. A DTA instrument

Thermal Analysis of Textiles and Fibers. DOI: https://doi.org/10.1016/B978-0-08-100572-9.00003-3 © 2020 Elsevier Ltd. All rights reserved.

18

Thermal Analysis of Textiles and Fibers

Figure 3.1 Schematic representation of a DTA instrument (S, sample; R, reference; TS, sample temperature; TR, reference temperature) holder to keep the two holders at the same temperature. The DTA curve is TS 2 TR 5 f(Tblock), where Tblock is the block (furnace) temperature.

consists of a furnace, which heats the sample and the reference contained in separate holders (see Fig. 3.1). Sensors are attached to the sample and the reference for measuring their temperature. The sensors most often are thermocouples or resistance thermometers. The signals from the sensors are used to determine the temperature difference between the sample and the reference (ΔT 5 TS 2 TR). Then ΔT is displayed as a function of the block (furnace) temperature (Tblock). Here TS is the sample temperature, and TR is the reference temperature. In earlier types of heat-flux DSCs the ΔT 5 f(Tblock) curve was the DSC curve. Obviously, in the temperature range without any heat-absorbing or heat-releasing process, ΔT will be close to zero (more precisely ΔT will have a small value proportional to the specific heat capacity change of the sample). On the other hand, if the sample melts during the DSC run (this is an endothermic process), ΔT 5 TS 2 TR will be negative. Let us suppose that we heat the holder containing the sample and the reference. The sample and the reference temperature will be going up parallel when no transition takes place, but the sample temperature will lag somewhat behind the reference temperature due to the heat capacity of the sample. However, when the sample starts melting, the sample temperature will stop at the melting point until the whole sample is melted (this is the criterion of equilibrium thermodynamics). At the same time the reference temperature keeps going up. DTA (differential thermal analysis) is somewhat different than DSC. In DTA the sample holder is directly filled with the sample. Often the reference holder contains some reference material. This reference material in most cases is aluminum oxide. Before the electronic revolution, DTA instruments worked with very large sample masses (usually grams). With such large sample masses, the response of the instrument is strongly influenced by heat-transport effects through the sample. Since thermal analysis measurements are primarily used for polymeric materials, thermal conductivity became a significant factor during the measurements. As a consequence, the packing and the amount of the sample also influence the quality of the results. In DSC instruments, there is no direct contact between the sample and the sensors: the sample is packed in a metal sample pan (the material of the sample pan in most cases is aluminum or gold, gold is for hightemperature measurements). Thus the sensors can be kept clean and will not get

Differential scanning calorimetry (DSC) in fiber research

19

Figure 3.2 (A) A schematic diagram for operation of a power-compensation DSC and (B) sample and reference holders of a power-compensation DSC. The DSC curve is the differential power versus the average temperature of the sample and reference (this is the same as the program temperature). Source: From Wunderlich B., 1990. Thermal Analysis. Academic Press, Boston, MA. Reprinted with permission from Academic Press.

G

contaminated with the molten sample. The mass of the sample in DSC is much smaller than that of in earlier DTAs, usually 1
10 mg (nanograms in fast-scan DSCs). In addition, in a DSC the physical form of the sample often is a thin layer (film). Thus the heattransport effects are minimized, and ΔT becomes almost proportional to the heat capacity of the sample. This type of DSC is called heat-flux DSC. Instrument manufacturers marketing heat-flux DSCs are TA Instruments, PerkinElmer, Mettler Toledo, Netzsch, Hitachi, Setaram, Shimadzu, Instrument Specialists, etc. (the DSC of Setaram is slightly different than that of the other companies, because it uses the Tian
Calvet principle). There is another type of DSC called power-compensation DSC (PC-DSC) (Fig. 3.2A and B). This instrument is manufactured by PerkinElmer only. In addition, the only commercial fast-scan DSC (the Flash2 DSC from Mettler Toledo) also works on the powercompensation principle. In the power-compensation DSC, there are two separate holders, one for the sample, the other one for the reference, but both are heated with individual heaters (see Fig. 3.2B). They both have separate platinum resistance thermometer sensors. An average amplifier heats both holders by the same current, and the signals from the sensors of the holders (the temperatures of the sample holder and the reference holder) are compared. In these DSCs the sample holder contains the sample in a sample pan, and the reference holder contains an empty sample pan. In simple heating experiments without a

20

Thermal Analysis of Textiles and Fibers

thermal transition, the sample holder is colder than the reference holder because of its higher heat capacity (it contains the sample in addition to the sample pan). A differential amplifier then puts additional power into the sample holder.

In both DSC types the DSC signal is more-or-less proportional to the specific heat capacity of the sample. Power-compensation DSC in this aspect is better: under optimum conditions, the specific heat capacity of a material can be measured to an accuracy of about 6 0.5%. In a DSC instrument, during heating, the sample temperature will lag behind the reference temperature (due to the higher heat capacity of the sample holder). During an endothermic process, the sample temperature stops at the melting point despite continuously increasing average or block temperature. From equilibrium thermodynamics, it is known that the material temperature stays constant at the melting point until the whole sample melted (at the end of wintertime when snow starts melting, the water temperature stays at 0 C until all the snow melts). When the melting is over, the sample temperature starts catching up the reference temperature, and the whole process is displayed as an endothermic peak. The area under the peak is proportional to the heat of fusion, and in general, the area under any peak (endothermic or exothermic) is proportional to the heat associated with the given process or transition. As mentioned earlier, the power-compensation DSC displays the DSC curve as the differential power versus the program (or average) temperature, while the heat-flux DSC displays the DSC curve as TS 2 TR versus the block (or T0) temperature. The DSC signal in both instruments (heat flux and power compensation) in steady state is more-or-less proportional to the specific heat capacity of the sample that is why both techniques are called DSC. The specific heat capacity in the power-compensation DSC is determined with a somewhat higher accuracy than in the heat-flux DSC (up to approximately 0.5% and 2%, accordingly). “Steady state” in systems theory is defined as a process in which the state variables determining the behavior of the system or the process are unchanged in time (Wikipedia; Gagniuc, 2017). For DSC measurements, this means no thermal transitions. But some change in the baseline is permitted, since the heat capacity of any material changes slowly with temperature. An endothermic transition in the DSC curve appears as a peak directed toward higher heat capacities, while an exothermic peak is displayed as a peak directed toward lower heat capacities. That is why in power-compensation DSC endo is up, while in heat-flux DSC endo is down. In modulated temperature DSC (MT-DSC), some modulation is added to the linear heating. This modulation can be sinusoidal, step-scan, and sawtooth. The following parameters are often used in DSC, and it will be useful for the readers to get familiar with these: 1. Thermal lag. This is the difference between the average sample temperature and the sensor temperature. The source of thermal lag is the thermal resistance between the sample holder and the bottom of the sample pan, the bottom of the sample pan and the sample (these two are called external thermal resistances), and also thermal resistance within the sample itself (called internal thermal resistance). The thermal resistance is the property

Differential scanning calorimetry (DSC) in fiber research

2.

3.

4.

5.

6.

7.

8.

21

of any material that slows down the flow of heat. Thermal resistance is the reciprocal of thermal conductance. It is desirable to minimize thermal resistance so that the real temperature of the sample is as close to the sensor temperature as possible. Temperature gradient. This quantity indicates the unequal distribution of the temperature in a material (it is the difference between the maximum and the minimum temperature in the sample). In DSC or other thermal analysis experiments, the temperature gradient within the sample depends on the sample mass, the heating rate, and the thermal diffusivity of the sample. It is desirable to minimize the temperature gradient in the sample in thermal analysis experiments. Heating rate. This quantity indicates how fast the sample is being heated in the instrument. The unit of the heating rate is  C/min (recently according to IUPAC K/min). The most frequently used heating rates in conventional commercial DSCs are between 1 and 40 C/min. For studies on reorganization of polymeric crystallites during heating, the heating rate can be as high as 500 C/min (PerkinElmer Hyper DSC). In the so-called fast-scan DSCs (FS-DSC, Flash1 DSC of Mettler Toledo), the heating rate can be as high as 2,000,000 C/min. Cooling rate. The unit of the cooling rate is  C/min (K/min). The most frequently used cooling rates are between 1 and 40 C/min. Sometimes quench cooling can be used for preparing amorphous samples of semicrystalline polymers, for example, poly(ethylene terephthalate) (PET), where the cooling rate can reach 2,000 C/min when using an efficient quenching agent (such as Freon). In the Flash1 fast-scan DSC of Mettler Toledo, the cooling rate can reach 200,000 C/min. Sample mass (in the United States still sample weight). Usually ranging between 1 and 10 mg, in fast-scan DSC it is nanograms. The allowed heating rate depends on the sample mass. The higher the heating rate, the lower the allowed maximum sample mass. The reason for this is that with increasing heating or cooling rates the thermal lag tends to increase. For a heating rate of 10 C/min, the sample mass should not exceed 10 mg. For a heating rate of 1 C/min, it can be B20 mg, while for 100 C/min around 1 mg. For very high heating rates of fast-scan DSCs, the sample mass is in the nanogram range. Atmosphere. The DSC measurements are carried out under the constant flow of some dry purge gas, usually high-purity nitrogen or argon. Helium would be very desirable because of its high thermal conductivity, but its high price limits its application. Sometimes nitrogen
helium mixtures are used due to their higher thermal conductivity, but they have lower price. The temperature calibration of a DSC instrument depends on the purge gas due to the difference in the thermal conductivity of the purge gases. Therefore if the purge gas is changed, the instrument needs to be recalibrated. Glass transition. It is a kinetic transition. This transition is always present in amorphous and semicrystalline polymers. Sometimes it cannot be easily detected on the DSC curves due to the extreme broadening of the transition and the presence of the oriented and rigid amorphous fractions (see Menczel and Wunderlich, 1981, 1986; Menczel and Jaffe, 2006, 2007), but dynamic mechanical analysis (DMA) measurements (and often modulated temperature thermomechanical analysis (MT-TMA) measurements) can always identify this transition. In the DSC curves, it appears as a heat capacity jump. It is characterized by the glass transition temperature (Tg, which is the temperature at half heat capacity increase) and the magnitude of heat capacity increase at the glass transition [ΔCp, J/(g  C) or J/(mol  C)] (see Fig. 3.3C). Melting. This is a first-order phase transition. At this transition, three-dimensional crystals become isotropic melt. In a DSC curve, it appears as an endothermic peak (as shown

Figure 3.3 (A) The full DSC curve (including the beginning and final isotherms) showing the melting of a high-purity low-molecular-mass sample (endotherm is down). T1 is the starting temperature of heating; at temperature T2, the steady-state conditions are achieved; at temperature T3, the melting starts due to minor amounts of impurities; between temperatures T4 and T5, the sample temperature is constant, this is a linear portion called the leading edge; between Tp (peak temperature, not shown in this diagram) and T6, the sample temperature increases again catching up the reference temperature; at temperature T6, steady state exists again; at temperature T7, the heating stops, and the temperature is kept constant at final isotherm. (B) The melting portion of the DSC curve from part (A) (without the beginning and final isotherms) showing the melting of a high-purity low-molecular-mass sample (in this case, indium, endotherm is down). At temperature T3, the melting starts due to minor amounts of impurities (the transition between T3 and T4 is curved); between temperatures T4 and T5, the sample temperature is constant, this “leading edge” is a straight line; the sample temperature catches up the reference temperature between Tp and T6; at temperature T6, steady state exists again. (C) The glass transition in a DSC curve and its calculation. The unfreezing of the sample starts at temperature Tb; then the heat capacity increase (ΔCp) accelerates, and it is completed at temperature Te; T1 and T2 are the extrapolated starting and end temperatures of the glass transition; Tg (midpoint) is the glass transition temperature. Source: From Menczel, J.D., Prime, R.B., 2009. Thermal Analysis of Polymers, Fundamentals and Applications. John Wiley and Sons. Reprinted with permission from Wiley.

Differential scanning calorimetry (DSC) in fiber research

23

in Fig. 3.3B). It is characterized by the heat of fusion (ΔHf) and the melting point (Tm). As will be seen later, for high-purity low-molecular-mass substances, the melting point is the extrapolated onset of the melting peak, while for polymers it is the last and the highest temperature point of the melting endotherm. For semicrystalline polymers, Tmp is also often given, this is the peak temperature of melting. Tmp does not have physical meaning, it is the temperature point where the rate of melting is the highest during the heating (for polymers only). It should not be used as the melting point (unfortunately, in many publications Tmp is reported instead of Tm). Why was the last point of the melting endotherm selected as the melting point? Because in polarization optical microscopy experiments this is the temperature at which the birefringence disappears during heating, so it is very convenient for the comparison of DSC and optical microscopy characteristic temperatures. 9. Melt crystallization. It is also a first-order phase transition where an isotropic melt becomes three-dimensional crystals. This transition takes place during cooling of a melt. For polymers, it is characterized by the heat of crystallization (ΔHc), the starting temperature of crystallization (Tco), and the peak temperature of crystallization (Tcp) (see Fig. 3.29B). Similar to Tmp, Tcp does not have a physical meaning, it is simply the maximum rate of crystallization during cooling. The heat of crystallization is not determined in most cases: the low-temperature baseline (i.e., the baseline below the crystallization transition) very often is not linear, so the sigmoidal baseline cannot be applied (see later in this chapter). 10. Cold crystallization. If the melt of a crystallizable polymer is cooled extremely fast (quenched), it may not be able to crystallize due to kinetic reasons, and it turns into a glass in a certain temperature range. When this (amorphous) glassy sample is heated in a DSC, the segmental motion starts at the glass transition temperature, and the sample crystallizes. This process is called cold crystallization because it takes place at temperatures much lower than the melt crystallization temperature (see Fig. 3.29A).

3.2

Calibration of differential scanning calorimetries

DSC instruments, like other thermal analysis instruments, need to be calibrated before the actual experiments. For DSCs, there are three different calibrations: temperature, energy (heat), and heat capacity.

3.2.1 Temperature calibration For accurate measurements the temperature should be calibrated separately for heating and cooling experiments.

3.2.1.1 Temperature calibration on heating Traditionally, temperature calibration of DSCs is done on heating. Most often, high-purity metal standards (In, Sn, and Pb) and spectroscopic purity water are used for this calibration. Fig. 3.3A and B shows the melting curve of high-purity indium standard. It can be seen that at around 156.6 C, an endothermic transition

24

Thermal Analysis of Textiles and Fibers

begins. This endothermic peak is the melting of indium. When we take a look at Fig. 3.3A (the full DSC curve), it can be seen that the DSC procedure starts with a sudden jump in the DSC signal (range T1
T2, called the starting transient). The magnitude of this jump is proportional to the specific heat capacity, and the mass of the sample, as well as the heating rate. This transient is used to measure the heat capacity of materials. A horizontal baseline then follows (temperature ranges from T2 to T3). In reality, this baseline is not absolutely horizontal but slightly sloped due to the increasing heat capacity of the sample. At temperature T3 the melting process starts. As can be seen in Fig. 3.3A and B, there is a small curved region between T3 and T4, this curvature exists due to the presence of minor amounts of impurities in the sample. For absolutely pure substances the start of the melting is a sudden break at the melting point from the horizontal baseline to the sloped “leading edge” of the peak, but as previously mentioned, minor amounts of impurities broaden the start of the melting, so the rise of the DSC signal at the start of the melting becomes a smooth, continuous process. Then a sloped straight line follows (range T4
T5), which is called “the leading edge” of the peak. It is linear, because in this temperature range, the sample temperature does not change for pure low-molecular-mass substances. According to equilibrium thermodynamics, the sample temperature during the melting process does not change until the whole sample is melted. The leading edge corresponds to this process. Then again, we have a small curved section at the top of the peak that corresponds to the end of the melting process. So, by now it is clear that the two small curved regions (T3
T4 and the peak maximum area) play no significant role in the analysis of the melting, therefore we will ignore them. Regions T2
T3 (the horizontal line, which is the premelting baseline) and the leading edge (T4
T5) are important. Remember that the leading edge is the part of the process where the sample temperature is constant at the melting point. So, we can summarize the melting as follows: when we heat the sample, no thermal event takes place in the temperature range T2
T3. With very high magnification, this part of the curve simply shows the temperature dependence of the specific heat capacity of the indium crystals. At T3 the sample starts melting, but this is premelting, taking place due to small amounts of impurities. In temperature range T4
T5 the sample temperature is constant, and essentially the melting is completed by temperature Tp. But at Tp the reference temperature is way up compared to the sample temperature. In the temperature range (Tp
T6) the sample temperature is catching up the reference temperature, and at temperature T6 we have steady state again. So, for the temperature calibration (determination of the melting point of indium), we need the temperature regions T2
T3 and T4
T5. The point of intersection of the extrapolated T2
T3 and T4
T5 lines is the melting point (this would be the melting point without any impurities in the sample). The melting point in Fig. 3.3B is designated as Tm. For the temperature calibration needed for heating experiments, we require the melting of at least two standards. Let us say, these are high-purity water and indium. Then, all the instrument operator has to do is to enter the two experimental and corresponding theoretical melting points into the instrument software. We need to remind

Differential scanning calorimetry (DSC) in fiber research

25

the readers that entering these values into the instrument does not cause any hardware changes. The later experiments will go on as before, but the obtained temperatures will be recalculated to real temperatures using the calibrating melting points. Of course, if the operator needs to check the linearity of the temperature scale of the instruments, more than two calibration standards are needed. It is recommended that a newly purchased instrument is run to determine the linearity of the temperature scale. If the temperature scale is linear, determination of the melting point of two standards will be sufficient later. If the temperature scale is not linear, three to four calibration standards are needed, and some mathematical function must be fitted to the experimental melting points. But all high-quality commercial DSCs must have linear temperature scale. If not, the readers are encouraged to call service. The following substances are recommended for the heating calibration of DSCs: G

G

G

For the negative temperature region (below 0 C): decane, Tm 5 229.2 C; octane, Tm 5 257.0 C. In the middle temperature region, spectroscopic purity water (melting point is 0.00 C). Here we need to remind that the software of the Q100, Q200, Q1000 and Q2000 TA Instruments DSCs will not work if 0.00 C is entered into the calibration table (this is due to some mathematical problems in the software). Therefore, it is recommended to enter 0.01 C as the literature melting point. In the higher temperature region, minimum 6
9 high-purity metals (6
9 means 99.9999% purity): indium (melting point is 156.60 C); tin (melting point is 231.90 C); and lead (327.47 C). Zinc does not always have a straight leading edge (because it alloys slightly with aluminum and the pan material is most often aluminum), therefore it is not recommended to use as a calibration standard.

There are two more standards that could be used as melting point standards, but they should not be used. These are mercury (Tm 5 238.8 C) and gallium (Tm 5 29.8 C). Mercury would be a desirable standard because there are no other very high-purity standards in the negative temperature region. Gallium also would be necessary, because it has a melting point around room temperature where no other metal melts. However, both of these metals are toxic, they both alloy with aluminum (the material of the sample pan), therefore they are not recommended as melting point standards. Except indium, most calibrating metals alloy with aluminum slightly when keeping them in the state of melt for sufficiently long time. This is especially true for zinc and tin. Therefore it is not recommended to use and remelt the same sample for many times. A tin sample can be used five to six times for melting, but for zinc the leading edge of melting loses the linearity after the first melting. This should be taken into account during the calibration. In this book, we are using the Celsius scale because most authors in this book are from the United States, and the United States did not catch up with the rest of the world yet with the temperature scales: in the everyday life the Fahrenheit scale is still used, and the Celsius scale is still in use in the scientific measurements. For recalculating, these temperature values to the thermodynamic temperature scale (K), 273.15 should be added to the degrees Celsius values.

26

Thermal Analysis of Textiles and Fibers

N-(4-n-Octyloxy-2-hydroxybenzal)-4'-n-butylaniline Transition

Tm (°C)

ΔHf (J/g)

Cr→SC

39.2

78.2

SC→N

69.4

2.57

N→I

89.2

4.6

Figure 3.4 The LC-1 liquid crystal for cooling calibration of DSCs (Menczel and Leslie, 1990, 1993; Menczel, 1994, 1997). Source: From Menczel, J.D., Leslie, T.M., 1993. J. Thermal Anal. 40, 957. Reprinted with permission from Springer.

3.2.1.2 Temperature calibration on cooling Temperature calibration on cooling could not be performed for a long time because the crystallization temperature of common substances differs considerably from the melting point, and it does not have a constant value. The reason for this difference is the so-called supercooling. Supercooling is not constant hence changes with the present foreign surfaces (i.e., in DSC experiments, the material of the sample pan) due to nucleation, the cooling rate, the purity of the sample, etc. It was in 1990 when Menczel and Leslie (Menczel and Leslie, 1990, 1993; Menczel, 1994, 1997) realized that the nematic-to-isotropic and the cholesteric-to-isotropic transitions of low-molecular-mass liquid crystals (LCs) do not have supercooling, the transition temperature of these substances on heating is identical to the transition temperature on cooling. So, the nematic-to-isotropic or cholesteric-to-isotropic transition can be used for the temperature calibration of DSCs on cooling. In the abovestated publications, the mentioned authors observed that the following transitions are without supercooling in addition to the isotropic-to-nematic or isotropic-to-cholesteric transitions: SJ !SI , SI !SC , SA !Ch and SC !N, so all these transitions can be used for cooling calibration of DSCs (here S means smectic, N means nematic, Ch means cholesteric, and I means isotropic). On this basis, Menczel and Leslie (1990, 1993) proposed that any LC!LC transition is without supercooling, so they all can be used for temperature calibration of DSCs and other analytical techniques on cooling. The first LC used for the cooling calibration was named LC-1 [N-(4-octyloxy-2-hydroxybenzal)
40 -butylaniline]. This LC has a crystal!smectic C transition at 38.9 C, a smectic C!nematic transition at 69.0 C, and a nematic!isotropic transition at 88.5 C (see Fig. 3.4). Menczel and Leslie (1990, 1993) observed the absence of supercooling for LCs if they had purity of at least 99.6
99.7 mol.%. These authors used two other commercial LCs for the cooling calibration (unfortunately, not available anymore commercially): LCs CE-3 and

Differential scanning calorimetry (DSC) in fiber research

27

(A)

(+) -4-n-Hexyloxyphenyl-4'-(2"-methylbutyl)-biphenyl-4-carboxylate Transition

Tm (°C)

Cr→SC*

66.0

ΔHf (J/g) 51.2

Sc*→Ch

78.8

3.35

Ch→I

163.5

2.62

Transition

Tm (°C)

ΔHf (J/g)

(B)

Cr→SJ*

55.8

54.6

SJ*→SI*

63.9

0.15

SI*→SC*

69.2

4.10

SC*→SA*

84.0

Not detectable by DSC

SA*→Ch

134.8

5.7

Ch→I

140.7

3.4

Figure 3.5 (A) Liquid crystal CE-3 for cooling calibration of DSCs (Menczel and Leslie, 1990, 1993; Menczel, 1994, 1997). (B) Liquid crystal CE-8 for cooling calibration of DSCs (Menczel and Leslie, 1990, 1993; Menczel, 1994, 1997). Source: From Menczel, J.D., Leslie, T.M., 1993. J. Thermal Anal. 40, 957. Reprinted with permission from Springer.

CE-8 of E. Merck [(1)-4-n-hexyloxyphenyl-2-hydroxybenzal)
40 -butylaniline] and [(1)-(4-(20 -methylbutyl)phenyl-40 n-octylbiphenyl-4-carboxylate], respectively (see EM Chemicals Technical Brochure EM-321S-84) (Fig. 3.5A and B). A precondition of using a LC for cooling calibration is the thermal stability of the standard, because these are organic substances. Since the transition temperatures are very sensitive to purity (degradation products may be present in the samples), the thermal stability can be checked if the standard is cyclically heated and cooled, and the transition temperatures are determined in every cycle. The results of such an experiment are shown in Fig. 3.6. Here it was observed that the transition temperatures of CE-8 did not change in four consecutive heating
cooling cycles, so the standard is thermally stable and can safely be used for cooling calibration experiments. For a cooling calibration the DSC instrument needs to be calibrated on heating at a specific heating rate using high-purity metal standards. Then a LC substance is run, and its transition temperatures should be determined on heating at the heating rate at which the instrument was calibrated by metals, and then the samples are run on cooling at the cooling rate of interest. If the transition temperature

28

Thermal Analysis of Textiles and Fibers

Figure 3.6 The DSC heating curves of liquid crystal CE-8 during cyclic heating
cooling experiments (endotherm is down) indicating the first and fourth heating. Heating and cooling rates are 2 C/min. The transition temperatures did not change from the first to the fourth heating, so the liquid crystal is thermally stable (Menczel, unpublished results). TA Instruments 2200/2910 DSC.

of N!I, Ch!I, or some LC!LC transitions are designated as Th, and the transition temperature measured on cooling is designated as Tc, then ΔT 5 Th 2 Tc will be a correction factor that needs to be applied to the transition temperature of the sample measured on cooling to obtain the correct temperature. Figs. 3.7 and 3.8 show the heating and cooling curves of LCs CE-3 and CE-8 recorded at the rate of heating and cooling at 2 C/min. It can be easily seen that the I!N, and all LC!LC transitions occur at the same temperature on heating and cooling, but the lowest LC to crystal transition has considerable supercooling. So, the procedure of the cooling calibration is as follows: 1. The instrument needs to be calibrated on heating at the desired heating rate (e.g., 10 C/min) with high-purity water and metal standards (In, Sn, and Pb). 2. Appropriate cooling calibration LC standards should be selected. Easily available LCs are listed in ASTM E2069 standard. The transition temperatures of the LCs should be determined on heating and cooling at desired heating and cooling rates. The transition temperature on heating and cooling is defined as the extrapolated onset of the transition. 3. The following temperature correction factor should be determined from the heating and cooling experiment:

ΔT 5 T h 2 T c

(3.1)

This correction should be applied to the cooling temperatures determined for any samples. In the equation, Th is the transition temperature determined from the heating experiment, and Tc is the transition temperature determined from the cooling run, ΔT is a correction factor. When ΔT is determined from several LC!LC

Figure 3.7 Heating and cooling DSC curves of liquid crystal LC-1 used for calibration on cooling. Heating rate 5 2 C/min, cooling rate 5 2 C/min (endotherm is down). The transition peaks from left to right are as follows: crystal!SC (Tm 5 39.2 C), SC!N (Tm 5 69.4 C), and N!I (Tm 5 89.2 C). It can be seen that the SC!N and N!I transitions occur at the same temperature in the heating and cooling runs, while the SC!crystal transition has a 20 C
25 C supercooling during cooling. TA Instruments 2200/2910 DSC. Source: From Menczel, J.D., 1997. J. Therm. Anal. 49, 193. Reprinted with permission from Springer.

Figure 3.8 Heating and cooling DSC curves of liquid crystal CE-8 used for calibration on cooling. Heating rate 5 2 C/min, cooling rate 5 2 C/min. The transition peaks from left to right are as follows: crystal!SJ (Tm 5 55.8 C), SJ !SI (Tm 5 63.9 C), SI !SC (Tm 5 69.2 C), SA !Ch (Tm 5 134.8 C), and Ch!I (Tm 5 140.7 C). The SC !SA transition could not be detected by DSC but was seen by polarization optical microscopy. All the transitions occur at the same temperature in the heating and cooling runs except the SJ !crystal transition, which shows a significant supercooling in the cooling run. TA Instruments 2200/2910 DSC. Source: From Menczel, J.D., 1997. J. Therm. Anal. 49, 193. Reprinted with permission from Springer.

30

Thermal Analysis of Textiles and Fibers

and/or LC!I (where LC 5 liquid crystal, I is the isotropic phase) transitions, it must be averaged to obtain an effective correction factor. The detailed procedure of performing the cooling calibration is described in ASTM Standard E2069-00 (2012).

3.2.2 Energy (enthalpy) calibration The enthalpy calibration is done simultaneously with the temperature calibration on heating. High-purity indium, tin, or lead are used since their heat of fusion is known to a high degree of accuracy. It is not really convenient to use water for this calibration. For temperature calibration, water is appropriate, but the high rate of evaporation makes the posttransition baseline nonlinear when using standard DSC pans. Water can be used as a calibration standard for energy calibration running it in hermetically sealed sample pans. Experience shows that the sample mass in this case must be smaller than 2 mg, otherwise the pans may open due to internal pressure. It should be mentioned that the heat of fusion of the mentioned metals has some scattering in the literature: ΔHf for indium is between 28.39 and 28.66 J/g, for tin between 58.97 and 60.22 J/g, and for lead, between 22.997 and 23.02 J/g (Haynes, 2014; Kaye & Laby Online, 2005; Coursey et al., 2014; Cottrell, 1954; Emsley, 2011; Aston et al., 1943; Ancsin, 1985; Archer and Rudtsch, 2003; Schumm et al., 1973). The recommended values for the heat of fusion of these metals are the following: for In, 28.66, Sn 60.22, and Pb 23.02. So, for the energy (and temperature calibration on heating) calibration the following procedure should be carried out: A relatively large piece of the metal standard should be cut out. Then a thin piece should be cut from the large piece so that the newly cut out piece does not have any surface being in contact with air previously. Any metal oxidizes, and the metal sometimes forms a eutectic mixture with its oxide leading to a melting point lower than the melting point of the metal itself. This would definitely falsify the calibration results. The melting and heat of fusion calibration software should be reset in the DSC instrument. The standard should be crimped in a standard aluminum DSC pan (water should be encapsulated in a hermetically sealed pan) and melted to increase the surface area of the standard piece. Then the standard should be cooled and remelted (second heating), and the results obtained from the second heating should be used for the calibration purposes. The melting point and the heat of fusion should be determined. The melting point should be taken as the extrapolated onset of the melting peak (the point of intersection of the premelting baseline and the leading edge). In addition, the heat of fusion should be determined using the straight baseline for the melting transition for metals. When determining the heat of fusion for water, the sigmoidal basene should be used because of the large difference in specific heat capacities of ice and water. This experiment should be repeated at least once and accepted when the two consecutive results are within 0.1% from each other for the energy, and the melting point results between two runs are less than 0.1 C. Then, the newly obtained calibration data should be entered the instrument software. The calibration standards should be run again in the calibrated DSC instrument, and the newly obtained results should be the same as the literature values for water, In, Sn, and Pb. By now, the instrument is calibrated and can be used for quantitative experiments.

Differential scanning calorimetry (DSC) in fiber research

31

3.2.3 Specific heat capacity calibration (heat flow calibration) Heat capacity is a function of state in thermodynamics. Heat capacity shows the amount of heat necessary to raise the sample temperature by 1 C. Its dimension is J/K. If the amount of the sample is 1 g or 1 mol, this quantity is called specific heat capacity, the dimension of which is J/(g K) or J/(mol K). In thermodynamics the following two heat capacity variants are used: G

G

Cv is the heat capacity at constant volume. Cp is the heat capacity at constant pressure.

The mathematical interdependence between Cp and Cv is as follows: Cp 5

dQ 5 Cv 1 dT 

Cp 5 CV 1 T

Cp 5 CV 1

@p @T

TVγ 2 βT

    @U @V 1p @V T @T p  v

@V @T

(3.2)

 (3.3) p

(3.4)


where γ  ð1=VÞ @V=@p
p is the volumetric coefficient of thermal expansion, and β T  2 ð1=VÞ @V=@p T is the isothermal compressibility. Polymers (including synthetic fibers and films) can exist only in solid or liquid states. Also, it is impossible to keep solid and liquid samples at a constant volume during heating (enormous pressures would be needed for this); therefore direct measurement of CV is impossible. Thus any DSC measurement leads to the determination of Cp. The DSC signal is more-or-less proportional to the heat capacity of the sample. This proportionality is better in power-compensation DSCs than in heatflux DSCs. Wunderlich and his research group (see, e.g., Gaur and Wunderlich, 1980, 1981, 1982; Varma-Nair and Wunderlich, 1990) collected many experimental data on heat capacity of various polymers, and these data were entered the ATHAS Databank. Lately this data bank belongs to Springer and can be accessed online at https://www.google.com/#q 5 Springer%2C 1 heat 1 capacity 1 of 1 polymers. The major advantage of DSC is the speed at which the heat capacity can be determined. The temperature dependence of the heat capacity of a material can be measured in several hours, and this is much faster than a comparative measurement using adiabatic calorimetry. Careful heat capacity DSC measurements can reach or exceed 6 1% accuracy of the results. The heat capacity in the absence of chemical processes increases with temperature, so the heat capacity of the liquid is always higher than the heat capacity of the glass or the crystal. Also, as was proven by Wunderlich, the heat capacity of a polymeric glass is close or identical to the heat capacity of the crystals of the same polymer.

32

Thermal Analysis of Textiles and Fibers

Figure 3.9 Overlay of the three DSC runs for determination of the heat capacity. (PerkinElmer DSC7 instrument). The shaded area is proportional to the heat input needed to heat the sample from T1 to T2. “Sapphire run” is the calibration run with a sapphire disk of known mass. “Sample run” is the actual run of the sample. “Baseline” is the run with two empty pans. Source: From Menczel, J.D., Prime, R.B., 2009. Thermal Analysis of Polymers, Fundamentals and Applications. John Wiley and Sons. Reprinted with permission from Wiley.

3.2.3.1 Calibration and determination of heat capacity with conventional differential scanning calorimetry Conventional DSC can be used to measure the heat capacity of materials. Fig. 3.9 shows three overlaid DSC runs: a blank run (also called a baseline, measured with two empty pans), a calibration run, and a sample run. In all of these runs the DSC cell is held isothermally for a short time (usually 2
5 minutes) at temperature T1 (called the starting temperature of the run) then heated at a linear rate. After the heating starts the DSC signal shifts, which continues until steady-state conditions are achieved. At temperature T2 (the final temperature of heating) the heating stops, and the sample is held isothermally at this temperature. When the heating stops, there is another signal shift (from the steady-state condition to the isothermal baseline). The area encompassed by the sample and the blank run curves (the shaded area in Fig. 3.9) is proportional to the heat input into the sample. Therefore an average heat capacity of the sample for the temperature range T1
T2 can be calculated using the following relationship: Cp 5

Q T2 2 T1

(3.5)

This procedure would be too lengthy to determine the temperature dependence of the heat capacity of a material: a separate run would be needed for T2 2 T1 5 1 C in order to determine the specific heat capacity at every 1 C.

Differential scanning calorimetry (DSC) in fiber research

33

Figure 3.10 The temperature dependence of specific heat capacity of polystyrene calculated from the data shown in Fig. 3.9. Perkin-Elmer DSC7, 10 C/min heating rate. Source: From Menczel, J.D., Prime, R.B., 2009. Thermal Analysis of Polymers, Fundamentals and Applications. John Wiley and Sons. Reprinted with permission from Wiley.

However, the curves from T1-to-scanning mode and scanning mode-to-T2 have similar shape; therefore the hs-bl (for sample and baseline) amplitude differences are proportional to the heat capacity at every temperature, and there is no need to determine the hs-bl area (shaded area in Fig. 3.9). Instead, the hs-bl amplitude differences should be determined. So, if the instrument is calibrated with a calibration standard for which the heat capacity is well known at all temperatures, the heat capacity of the sample can be determined at any temperature. For this the hsap-bl amplitude differences must be used (see Fig. 3.9, “Sapphire run” and “Baseline” curves). The standard is usually a sapphire (crystalline Al2O3) disk, which is readily available from the instrumental companies. For low-temperature heat capacity measurements, benzoic acid or copper is often used as a standard, because sapphire’s heat capacity changes fast with temperature in the subambient region. It should be mentioned that not all instrumental companies have standard software for heat capacity measurements: often the heat capacity software should be separately purchased. In such cases the heat capacity can be calculated using the procedure described below. The Cp measurements of Fig. 3.9 were made on a Perkin-Elmer DSC7 instrument. The shaded area in Fig. 3.9 is proportional to the heat necessary to raise the temperature of the sample from T1 to T2. The specific heat capacity shown in Fig. 3.10 was calculated from the three runs of Fig. 3.9. Bares and Wunderlich (1973) claim that better Cp results can be obtained if the sample pan masses are matched (to at least 6 0.2%). This is strange because in this case the additivity of the heat capacity would be questioned, and this would work against the first law of thermodynamics. It is noteworthy that the software of the TA Instruments’ Q2000 and Discovery Q2500 DSCs is capable of correcting different pan masses. At the same time, we need to mention that no systematic study is

34

Thermal Analysis of Textiles and Fibers

available showing the influence of such a correction on the accuracy of the Cp determination. So, to calibrate the instrument in scanning mode for heat capacity, the following procedure should be applied: 1. The masses of four sample pans with lids (two pans for the baseline, one pan for the sapphire standard, and one pan for the sample) should be determined for an accuracy of at least 6 0.1% or better in order to ensure a better than 6 1% accuracy of the heat capacity the sample. The masses of the sample pans should not differ from each other more than 6 0.2%. 2. Two empty pans should be selected to run an instrumental baseline (zeroline). All run parameters should be identical in the baseline, sapphire, and sample runs. Select the temperature range for the heat capacity data. The first B10 C
20 C from this temperature range will be lost from the point of view of usable results. This is the temperature range needed to achieve steady-state conditions. The time in the starting and final isotherms should be around 5 minutes, the heating rate is preferably 10 or 20 C/min. When the cell temperature is lowered to the temperature of the starting isotherm, it is preferable to wait for 10 minutes, and then the program can be started (temperature equilibration). 3. A sapphire disk with a known weight (around 20
50 mg) should be put and crimped in the third pan, and the previous experiment should be repeated with the sapphire pan and one empty (reference) pan. 4. The amplitude differences between the sapphire and baseline should be determined at all temperatures of interest, as shown in Fig. 3.9. Then, the instrument constant should be determined using the following equation: KðCp Þ 5 Cp;sap ðTÞ

msapUq hsap-bl ðTÞ

(3.6)

where K(Cp) is the instrument constant for specific heat capacity determinations, Cp,sap(T) is the specific heat capacity of sapphire at temperature T, msap is the mass of the sapphire disk, q is the heating rate ( C/min), and hsap-bl(T) is the amplitude difference between the sapphire and baseline amplitudes at temperature T. The heat capacity constant K(Cp) of the power-compensation DSCs should not show any dependence on temperature, but in heat-flux DSCs the “Cp constant” is not really constant. Therefore in heat-flux DSCs, it is recommended to determine K(Cp) for every 10 C. 5. The sample should be weighed in the fourth pan and run with the same conditions as the sapphire standard and the baseline. 6. The specific heat capacity of the sample can be calculated using the following equation: Cp;s ðTÞ 5 KðCp Þ

hs-bl ðTÞ msUq

(3.7)

where Cp,s(T) is the specific heat capacity of the sample at temperature T, hs-bl(T) is the amplitude difference between the sample and baseline runs at temperature T, ms is the mass of the sample, and q is the heating rate (K/min or  C/min).

For accurate heat capacity determination, it is recommended to determine the heat capacity of at least three samples from the same material.

Differential scanning calorimetry (DSC) in fiber research

35

The accuracy of heat capacity determination with a conventional DSC can be better than 6 1% when the measurements are done carefully with a powercompensation DSC (Bares and Wunderlich, 1973). The instrumental baseline depends on the environmental temperature. Therefore for good reproducible results, all the runs should be made at constant environmental temperature (the room temperature deviation should be less than 6 2 C). This requirement is not that critical these days due to usage of powerful cooling units. Naturally, all the runs (sample, baseline, and calibration) should be performed with identical parameters (identical starting isotherm in both temperature and time, identical heating rate, and identical final isotherm in both temperature and time, identical purge gas type and flow rate). The temperature range of the runs should not exceed 100 C. The following conditions are critical when determining the heat capacity: 1. If the sample is in molten state at temperature T2 (the final temperature of the run), care must be taken to avoid thermal degradation. 2. The isothermal baselines at temperatures T1 and T2 must be horizontal. If any sloping is observed, erroneous heat losses may take place (e.g., the sample holder lid may be tilted), or some chemical or physical process may be proceeding. This leads to unacceptable results. If the baseline at T2 is not horizontal, but the T1-baseline is horizontal, some chemical process (e.g., degradation) may be proceeding at the final temperature of the run. If both the low-temperature and the high-temperature baselines are sloped, the sample holder cover may be tilted. 3. The sample size should be 10
20 mg. The preferred heating rate is 10 or 20 C/min. The sample size in the fast-scan DSCs (e.g., Flash 2 DSC of Mettler Toledo or Schick’s fastscan DSC, see Schick and Mathot, 2016) is in the nanogram range. In these experiment special procedures are necessary to determine the sample mass. 4. The sample often must be conditioned before running. Many polymers (e.g., nylons) contain water. This water needs to be removed from the sample before the Cp determination, because its presence may falsify the results (nylons may contain water up to B8%), and accelerated degradation can be resulted at high temperatures. 5. The DSC sample holder must be dry. As an example, in a PerkinElmer DSC7, ice condensation must be prevented on the swingaway enclosure cover to ensure reproducible baseline. Therefore the purge gas must be dried before entering the cell, and dry conditions must be ensured in the dry box.

Sometimes the heat capacity determination must be carried out on cooling, especially for determining melt heat capacities. It is not easy to determine the temperature dependence of the melt heat capacity with traditional heating experiments, because the available temperature range of the melt is often limited by the onset of thermal degradation. The temperature dependence of the melt heat capacity is frequently needed, and this dependence should often be extrapolated to lower temperatures. Since supercooling is a common phenomenon in polymer crystallization, the temperature range available for heat capacity determination of the melt during cooling can be expanded when compared to heating runs. Naturally, the other two measurements (standard and baseline runs) and the temperature calibration of the DSC instrument will also have to be carried out on cooling.

36

Thermal Analysis of Textiles and Fibers

3.2.3.2 Measurement of heat capacity with modulated temperature differential scanning calorimetry in the quasiisothermal mode MT-DSC can make accurate heat capacity measurements in quasiisothermal mode (Jin et al., 1993; Boller et al., 1994; Thomas, 2006; Reading and Hourston, 2006; Ishikiriyama and Wunderlich, 1997), but the measurements are lengthy, since Cp is determined point by point. In these measurements the amplitude of the modulated heat flow is used for determining the reversing heat capacity (in the software of DSC 2920 this is called complex heat capacity). Therefore the need for baseline run is eliminated. In MT-DSC the following equation is used to determine the reversing heat capacity: Cp;rev 5 KðCp Þrev

HFlA HRA

(3.8)

where K(Cp)rev is the calibration constant for the reversing heat capacity, HRA is the heating rate amplitude, and HFlA is the heat flow amplitude. In these measurements the modulation must not be fast (the modulation period must be at least 120 seconds) to ensure that sufficient heat transfer exists between the sensor and the sample. The temperature modulation is preferentially 0.5 C
1.0 C, and a modulation period is $ 120 seconds as previously mentioned. In cases when the sample may release certain low-molecular-mass substances, hermetically sealed pans are used to prevent the sample holder contamination. When using hermetically sealed pans, the modulation period must be increased, because the contact area between the pan and the sensor is smaller than that for standard DSC pans. We will not discuss further the basics of DSC in detail, instead we send the interested readers to one of the thermal analysis textbooks (see, e.g., Turi, 1997; Menczel and Grebowicz, 2019, 2020, 2021; Menczel and Prime, 2009; Wunderlich, 1990; Reading and Hourston, 2006; Schick and Mathot, 2016) and will focus on special features of DSC measurements of fibers.

3.3

Differential scanning calorimetry of fibers

Melting experiments provide the most detailed information about the fiber structure. The melting point (Tm for polymers is the end temperature of the melting peak, see Fig. 3.15) determines the ultimate use temperature of the fiber, and it is necessary to establish the annealing conditions during manufacturing of the fiber. And these annealing conditions are needed to increase the fiber stability. We need to mention the appearance of multiple melting peaks for artificial fibers. Multiple melting is a frequent phenomenon in DSC curves of polymers (so, also of fibers), and it may have various sources, which is described as follows:

Differential scanning calorimetry (DSC) in fiber research

37

1. The fiber may have more than one component. In this case the melting of all components will be seen on the DSC curve (e.g., YONEX BG80LD badminton string composed of Vectran and Nylon). 2. Reorganization during melting. Wunderlich was the first to recognize that in semicrystalline polymers crystal perfection and recrystallization during melting frequently occur. The difference between these two is that crystal perfection (also called reorganization) takes place in the solid phase, that is, in crystals without melting. On the contrary, during recrystallization, the crystals melt and recrystallize in a more prefect crystal form. This process usually leads to a double-melting peak, since it shows the starting of endotherm (melting of the original crystals), an exotherm of the recrystallization process, and finally the melting of the recrystallized material. 3. The heat treatment at several discrete temperatures often gives multiple melting peaks (see, e.g., Varga et al., 1979 for low-density polyethylene (PE) films). 4. Small endotherms can often be observed on the DSC heating curves of semicrystalline polymers, including fibers (Berndt and Bossmann, 1976; Jaffe, 1978; Oswald et al., 1977; Valk et al., 1971; Valk, 1972; Wiesner, 1974, 1976a,b). Most of these peaks represent “annealing peaks”: if a fiber is kept for a certain time at a certain temperature beyond the glass transition temperature, small crystallites will develop, and their size and perfection will depend on the heat-treatment temperature. Therefore, the melting point of these crystallites will be determined by the temperature and time of the heat-treatment. This is the origin of the annealing peaks, sometimes called “middle endothermic peaks”. 5. One sample may contain various crystal forms of the same polymer. A good example is polypropylene (PP): crystallites of the β- and γ-forms of PP develop under various stresses, thus β- and γ-crystallites may appear as a result of stretching of the fiber or film (Varga, 1983). 6. Crystal-to-crystal transitions can sometimes be observed (Menczel et al., 1996; Menczel, 2020a). Poly(2-methylpentamethylene terephthalamide) (Nylon M5T) is an example for this when it is heated in the DSC (see “Amorphous A-R (as-received) curve in Fig. 3.11). Here the heat capacity jump at around 140 C corresponds to the glass transition of the amorphous polymer (heating rate is 5 C/min), then three exothermic peaks can be seen. The first one of these is the cold crystallization corresponding to the amorphous!crystal A transition. After that, the just-developed crystal melts and crystallizes into crystal form B (so this is an exotherm with an overlayed endotherm); then crystal form B melts, and the melt is crystallizing into crystal form C (third exotherm with some endotherm), and finally crystal form C melts (the highest temperature endotherm) into the smectic phase and then into isotropic phase. Crystal forms A and B are monotropic, while crystal form C and the smectic phase are enantiotropic. The sample containing the various crystal forms can be prepared by heat treatment of the as-spun fiber, and the DSC curves of these samples are also shown in Fig. 3.11. Thus the presence of multiple peaks is a complex event, and the number of peaks does not have to be constant, it may depend on the experimental conditions. The just-described as-spun Nylon M5T fiber may show fewer peaks if the heating rate is increased (Fig. 3.12): if the heating rate is raised to 40 C/min, there is simply no time for the amorphous!crystal A, crystal A!crystal B, and crystal B!crystal C transitions. In this case the amorphous sample cold crystallizes into crystal form C that will further melt into the smectic phase; the smectic phase will then melt into isotropic liquid.

It is not easy to prepare fiber samples for DSC measurements. These drawn fibers have large aspect ratio and are often of extremely small diameter (down to 10 nm). Also,

Figure 3.11 The heating DSC curves of amorphous Nylon M5T (as-spun fiber) and the other crystal forms of this polymer prepared by heat treatment. A-R stands for “as-received”. “No chain ext.” indicates the polymer prepared without a chain extender, thus having a relatively low molecular mass. Heating rate 5 5 C/min. TA Instruments Q2000 DSC. Source: From Menczel, J.D., 2020a, to be published in JTAC.

Figure 3.12 The heating rate dependence of the crystal!crystal transitions for amorphous Nylon M5T. TA Instruments 2200/2910 DSC. Source: From Menczel, J.D., Jaffe, M., Saw, C.K., Bruno, T.P., 1996. J. Therm. Anal. 49, 201. Reprinted with permission from Springer.

Differential scanning calorimetry (DSC) in fiber research

39

the small fiber pieces tend to “fly apart” when one tries to put them into a DSC pan. Therefore it is a problem to put a sufficient amount of fiber pieces into the DSC pan. As described in Chapter 1, Introduction, most commercial fibers are drawn, therefore they are oriented. The consequence of orientation is shrinkage during heating above the glass transition temperature. The effect of shrinkage will be seen on the melting curve of the fiber. Therefore if shrinkage is prevented during the heating, the melting curve will be different, and the melting point will be higher than for the samples where shrinkage is allowed. Thus two different DSC measurements can be performed (Mead and Porter, 1976; Miyagi and Wunderlich, 1972; Samuels, 1975a,b; Tashiro et al., 1980; Todoki and Kawaguchi, 1977a,b): 1. “Free-to-shrink” or “unconstrained” measurements. In this case the fiber is chopped into small pieces, which are then put into the DSC pan, the pan is crimped and put into the cell of the DSC instrument for a measurement. It is important to note that the fiber pieces can shrink during the measurement, because their ends are not fixed. It is also important that the measurements should be carried out in the so-called standard DSC pans (Fig. 3.13). In these pans a circular lid covers the fiber pieces. During the crimping process, the crimper pushes down the lid on the sample and folds the edges of the pan. So in this type of a pan, the lid is pushed down on the sample, so there is sufficient contact between the fiber pieces and the bottom of the pan (as often said, “There is good thermal contact.”). But the contact is not too tight, so the fiber pieces are free to shrink. If hermetically sealed DSC pans were used in these measurements, the lid would not push down the fiber pieces to the bottom of the pan, so the thermal contact would not be sufficient. In such cases the thermal lag would be excessive, and the DSC results would be distorted, therefore hermetically sealed DSC pans are not recommended for DSC measurements of fibers in free-to-shrink state. 2. “Constrained” or “fixed-length” measurements. Miller (1971) built a DTA instrument specifically designed for fiber measurements. In this instrument, long fiber pieces are wound in grooves on the outer surface of an aluminum cylinder. The ends of the fibers are fixed by tying them or gluing, thus the length of the fiber will not change during the heating. The aluminum cylinder sample holder also allowed for free-to-shrink measurements if short fiber pieces were put inside the hollow cylinder. This design was very innovative but could not be the solution for DSC measurements of fibers. And in the 1970s, DSC replaced DTA for measurements on plastics. The solution of carrying out constrained DSC measurements on drawn fibers is very simple. The fiber is wound on a thin stainless-steel circular piece in which two V-shaped cuts are made on the opposite sides of the plate (see Fig. 3.14), and the ends of the fibers are fixed somehow (tying together or glued), so the fiber will retain its original length during the measurement. To fix the ends of the fiber is not easy. A good but lengthy and difficult process is to tie the ends of the fiber. To fix the ends of the fiber by gluing is not always a good solution: the glue will display an endothermic or exothermic peak on heating making the analysis of the data difficult. A good solution to this problem is if the two ends of the fiber are sticking out of the pan before crimping: the crimping process will fix the ends of the fiber, because the pan edges are folded down on the lid. But even in this case, the ends of extremely thin fibers (B10 nm in diameter) may slip making partial shrinkage possible. It is also important to select an optimum thickness of the stainless-steel plate. By the experience of the authors, 0.12 mm is the optimum thickness: in this case the mass of the plate is approximately 25 mg. These plates can be cut reproducibly, so excessive sloping of the baseline

40

Thermal Analysis of Textiles and Fibers

Figure 3.13 The standard DSC sample pan: (A) A DSC standard pan, a standard lid, and a lid punched out of aluminum mesh (the pan and the standard lid are from Perkin-Elmer; the aluminum mesh lid was made by the first author). (B) A standard DSC aluminum pan crimped with a standard lid and a lid punched out of aluminum mesh (Perkin-Elmer pans). Source: From Menczel, J.D., Prime, R.B., 2009. Thermal Analysis of Polymers, Fundamentals and Applications. John Wiley and Sons. Reprinted with permission from Wiley.

Figure 3.14 The circular steel plate for measuring constrained melting of drawn fibers. Source: From Menczel, J.D., Prime, R.B., 2009. Thermal Analysis of Polymers, Fundamentals and Applications. John Wiley and Sons. Reprinted with permission from Wiley. can be avoided. If the plate is too thick, the instrumental baseline may be distorted, because a similar plate needs to be put in the reference side, and the mass difference between the sample side plate and the reference side plate sometimes may be too large. In

Differential scanning calorimetry (DSC) in fiber research

41

Figure 3.15 Determination of the characteristic temperatures and transition parameters (Tg, the glass transition temperature; ΔCp, the heat capacity jump at the glass transition; Tmp, the peak temperature of melting; Tm, the melting point; and ΔHf, the heat of fusion of a semicrystalline polymer). Note that the sigmoidal baseline is used to determine the heat of fusion. TA Instruments Q2000 DSC.

such cases the instrumental baseline would be sloped making evaluation the DSC curve more difficult. On the other hand, if the plate is too thin, it may be deformed during the windup of the fiber.

A number of researchers noted differences in the melting points and the general melting behavior between the melting in the free-to-shrink and constrained modes, such as the melting point in the constrained measurements is always higher than in the free-to-shrink measurements. This can be easily explained if we remember that Tm 5

ΔHf ΔSf

(3.9)

Fig. 3.15 shows how to determine the characteristic transition temperatures from a DSC run. In Eq. (3.9), Tm is the melting point, ΔHf is the heat of fusion, and ΔSf is the entropy change during the melting. Obviously during constrained melting, the entropy change will be much smaller than in the free-to-shrink measurement, because the polymer segments will not be able to relax immediately beyond the melting point, they will tend to stay oriented. So, if the heat of fusion is identical in the two types of measurements (and this is not always true), the melting point in the constrained measurements must be higher. This can be observed in the actual measurements. At the same time, we need to mention that the heat of fusion also depends on the extent of the restraint of the fiber. Smook and Pennings (1984) compared the heats of fusion values of drawn gel-spun PE in the two types of measurements and observed a B7% difference between the two (see Figs. 3.16
3.18). The authors explained these differences by the

42

Thermal Analysis of Textiles and Fibers

Figure 3.16 (A) DSC melting curves of ultrahigh-molecular-mass gel-drawn polyethylene fiber in free-to-shrink state. The fibers were hot-drawn at 148 C. The draw ratios are indicated at each curve. Perkin-Elmer DSC-2 instrument, heating rate 5 C/min. (B) DSC melting curves of ultrahigh-molecular-mass gel-drawn polyethylene fiber in constrained state. The fibers were hot-drawn at 148 C. The draw ratios are indicated at each curve. Perkin-Elmer DSC-2 instrument, heating rate 5 C/min. Source: From Smook, J., Pennings, J., 1984. Colloid Polym. Sci. 262, 712. Reprinted with permission from Steinkopff Verlag Darmstandt.

development of partially ordered melt in the fixed-length DSC measurements. This is the same effect that was described earlier with the decreased entropy change of melting in the constrained measurements. It is clear that the orientation developed in the melt in the constrained DSC measurements is not stable thermodynamically, and it relaxes slowly. Another important effect in the DSC measurements of fibers is superheating. Zachmann (Zachmann, 1965; Zachmann and Spellucci, 1966) noticed that the extended-chain crystals of polymers tend to superheat during melting. He gave a reasonable theoretical explanation for this superheating that around the extended chains (or in drawn fibers, around the oriented amorphous regions) the conformational freedom is limited. On the other hand, if the entropy change is mitigated by the decrease of stress (like in bulk amorphous semicrystalline polymers), no superheating should take place as it is usually observed. Jaffe and Wunderlich (1967) observed superheating in extended-chain selenium crystals (remember: selenium is a polymer!). Later, Miyagi and Wunderlich (1972) noticed that crystals, much less perfect than extended-chain equilibrium crystals, namely, oriented crystals in drawn PET fibers, can also superheat (Fig. 3.19).

Differential scanning calorimetry (DSC) in fiber research

43

Figure 3.17 Heats of fusion ΔH1 (Fig. 3.17A), ΔH2 (Fig. 3.17B), and ΔH3 (Fig. 3.17C) assigned to the melting peaks in the constrained melting DSC curves shown in Fig. 3.16B. Perkin-Elmer DSC-2 instrument, heating rate 5 C/min. Source: From Smook, J., Pennings, J., 1984. Colloid Polym. Sci. 262, 712. Reprinted with permission from Steinkopff Verlag Darmstandt.

Todoki and Kawaguchi (1977a,b) were the first to carry out detailed investigation on free-to-shrink and constrained-state fiber melting using drawn Nylon 6 fibers. They observed a double-melting curve in the free-to-shrink measurements and assigned the double melting to crystal perfection of the original crystals, simultaneous melting and

Figure 3.18 The heat of fusion of ultrahigh-molecular-mass polyethylene fiber as a function of draw ratio measured in free-to-shrink and constrained states. The average difference between the two techniques is 17 J/g. This difference was explained by the development of partially oriented melt in the constrained-state measurements. Source: From Smook, J., Pennings, J., 1984. Colloid Polym. Sci. 262, 712. Reprinted with permission from Steinkopff Verlag Darmstandt.

Figure 3.19 DTA melting peak temperatures of poly(ethylene terephthalate) tire yarn as a function of heating rate for different sample preparation techniques: FO (restrained) (constrained) measurements; FO (unrestrained): free-to-shrink measurements; F2: the FO sample was annealed for 24 h at 250 C in vacuum, free-to-shrink measurement. DuPont (TA Instruments) DTA, sample mass was 1
2 mg. Source: From Miyagi, A., Wunderlich, B., 1972. J. Polym. Sci., Polym. Phys. Ed. 10, 1401. Reprinted with permission from Wiley.

Differential scanning calorimetry (DSC) in fiber research

45

recrystallization of the perfected crystallites, and finally melting of the recrystallized crystals. At the same time the authors observed a single melting peak in the constrainedstate measurements at higher temperatures than in the free-to-shrink experiments. According to these authors, the melting point of the fibers in the constrained-state measurements increased linearly with the draw ratio, and the temperature of drawing did not affect the melting point. In these papers and in the paper published by Todoki (1985), the following three techniques were used to characterize the fibers (see Fig. 3.20): 1. Technique A: the fibers were kept in acetylene atmosphere and irradiated with X-rays. The goal of this procedure was to create a cross-linking network in the amorphous regions (X-rays will not cross-link the polymer chains in the crystals). Obviously, at these conditions, reorganization of the imperfect original crystallites is prevented, and one obtains the melting curve reflecting behavior of the original crystallites. In these experiments a single melting curve is observed with a peak temperature of about 190 C. 2. Technique B: measuring the DSC curve of the fiber in the free-to-shrink state by a conventional technique (i.e., the fiber is not treated with acetylene). As mentioned earlier, for drawn Nylon, a double-melting peak is obtained with melting peak temperatures of 218 C and 225 C. In this method the melting curve does not reflect the melting of the original crystallites, but it provides information about the various processes taking place during heating in the free-to-shrink state. 3. Technique C: constrained-state measurements. There is no fiber shrinkage in these measurements. The melting point is increased due to entropy effects as previously described. The melting point is higher than in the free-to-shrink experiments, and also the melting peak is broader than in the free-to-shrink measurements. Illers (1970) showed that, in addition to entropy effects, the melting point is also increased due to superheating effects when the heating rate is higher than 0.5 C/min!!! Also, the heat of fusion was slightly lower in fixed-length measurements in good agreement with later measurements of Smook and Pennings (1984). Today we know that the lower heat of fusion in the constrained measurements is the consequence of the formation of the oriented melt (“pseudo-nematic” state). The melting point increase was consistent with annealing and superheating experiments indicating that superheating will take place if the oriented small extended-chain crystallites are immobilized by constraint. The melting point increase in the constrained DSC measurements is a general phenomenon, and it is true for all semicrystalline fibers (see, e.g., Fig. 3.21).

In the free-to-shrink measurements the observation of the double-melting curve by Todoki and Kawaguchi (1977a,b) was in good agreement with earlier results: metastable polymer crystallites exhibit crystal perfection during heating or in annealing treatment. This crystal perfection process can have various forms, such as crystal perfection in the solid state or recrystallization during melting. Menczel (2020c) identified the process of reorganization during free-to-shrink melting of PET, Nylon 66, and PP (polypropylene) fibers with MT-DSC by obtaining an intense exotherm in the nonreversing heat flow. The reorganization in the nonreversing heat flow of MT-DSC was always found, even when the double-melting peak could not be observed in the total heat flow curve. These observations confirmed the data obtained earlier by Miyagi and Wunderlich (1972) (see later). Menczel (2020b) also found by MT-DSC that no reorganization takes place during constrained melting (Fig. 3.22A and B).

46

Thermal Analysis of Textiles and Fibers

Figure 3.20 Figure (A) DSC melting curves of nylon 6 yarn obtained with different sample preparation techniques. Technique A: a fiber whose amorphous portions were cross-linked by γ-irradiation in acetylene; Technique B: conventional DSC technique on as-received fibers in free-to-shrink state; Technique C: the melting curves of fiber in constrained state (the fiber was wound up on an aluminum plate). Endotherm is down. Figure (B) The melting points of nylon 6 yarns as a function of draw ratio and sample preparation technique. The techniques are the same as in Figure 3.20A. The temperatures indicated at each curve show the draw temperatures. Measurements on a Perkin-Elmer DSC-1B, heating rate 10 C/min. Source: From Todoki, M., Kawaguchi, T., 1977a. J. Polym. Sci., Polym. Phys. Ed. 15, 1067; Todoki, M., Kawaguchi, T., 1977b. J. Polym. Sci., Polym. Phys. Ed. 15, 1507. Reprinted with permission from Wiley.

Differential scanning calorimetry (DSC) in fiber research

47

Figure 3.21 DSC melting curves of polypropylene fiber in free-to-shrink and constrained states. Dashed line: constrained measurement; solid line: free-to-shrink measurement. Heating rate is 10 C/min. TA Instruments Q2000 DSC (Menczel, 2020c).

Melting experiments are the most useful ones for clarifying the fiber structure because in polymers the crystal size and perfection are sensitive to changes in thermal and mechanical history. Also, the melting point characterizes the ultimate use temperature of the fiber and gives some information about setting the annealing conditions for achieving increasing fiber stability. It needs to be mentioned that annealing at temperatures somewhat lower than the melting point, helps to raise the melting point. In nonoriented polymeric samples (extended-chain selenium), Jaffe and Wunderlich (1967) observed superheating. They gave a reasonable explanation for the superheating of extended-chain crystals including thermal conductivity and layering of the extended-chain macromolecular segments from the surfaces of the extended-chain crystals. Later Miyagi and Wunderlich (1972) found superheating of oriented crystallites in PET fibers in constrained-state measurements (see Fig. 3.19). It is clear that the crystallites in these fibers were much less perfect than the equilibrium extended-chain crystals, so the superheating did not take place due to heat conduction, but due to entropy restrictions of the tie molecules that connect the crystallites and due to mobility effects on the outer layers of the oriented crystallites. Miyagi and Wunderlich (1972) showed that reorganization of the crystallites is present in the free-to-shrink experiments. On the other hand, in constrained measurements of drawn (i.e., oriented) PET fibers superheating was observed at heating rates at or above 5 C/min. In the constrained measurements at a heating rate of 5 C/min, the melting curve was very narrow, and the melting was very fast. At lower and higher heating rates the melting peak broadened. It seems that simultaneous crystal perfection, recrystallization, and superheating are responsible for this phenomenon: at low heating rates crystal perfection and recrystallization

48

Thermal Analysis of Textiles and Fibers

0.0

0.5

Non-rev. heat flow (W/g)

Endotherm

Heat flow (W/g)

–0.5

–1.0

–0.2 –0.4

–0.4 –0.6

–0.6 –0.8

–1.0 250

–1.5 0

–50

Rev. heat flow (W/g)

0.0 –0.2 0.0

50

100

150

200

Temperature (°C) 0.5

0.0

–1.0

–1.5

–2.0 –50

0.2 –0.5 0.0 –1.0

–0.2

Rev. heat flow (W/g)

Non-rev. heat flow (W/g)

–0.5

Endotherm

Heat flow (W/g)

0.0

–1.5 –0.4

0

50

100

150

–2.0 200

Temperature (°C)

Figure 3.22 (A) The melting curves of polypropylene fiber recorded in the constrained state using MT-DSC. The nonreversing heat flow does not indicate any reorganization during melting. The curves from top to bottom: reversing heat flow, total heat flow, non-reversing heat flow. TA Instruments Q2000 DSC (Menczel, 2020c). (B). The melting curves of PET fiber recorded in the free-to-shrink state using MT-DSC. The nonreversing heat flow indicates significant reorganization during melting. TA Instruments Q2000 DSC (Menczel, 2020a).

dominate the melting, but there is no superheating. At high heating rates the crystal perfection and recrystallization during melting are suppressed, and superheating phenomena dominate the melting. The combination of these three effects leads to the narrowest melting curve at a heating rate of 5 C/min. The width of the melting curve is not clear. Miyagi and Wunderlich (1972) obtained very narrow melting

Differential scanning calorimetry (DSC) in fiber research

49

Figure 3.23 DSC curves indicating glass transition of polystyrene of narrow-molecular-mass distribution on heating of samples previously cooled at 0.1, 1, and 10 C/min. Heating rate: 10 C/min. TA Instruments Q2000 DSC. Source: From Menczel, J.D., Prime, R.B., 2009. Thermal Analysis of Polymers, Fundamentals and Applications. John Wiley and Sons. Reprinted with permission from Wiley.

curves in constrained melting, while Smook and Pennings (1984) and Menczel (2019) observed the narrow melting in free-to-shrink experiments rather then in constrained one. In the last experiments, the melting range was broader, but the melting point was higher. We need to talk about the glass transition in fibers. Sometimes it is not easy to notice the glass transition, or the heat capacity increase on the DSC curves may correspond to other physical or chemical phenomena, but not the glass transition. In such cases the phenomenon of enthalpy relaxation may help to decide whether we have a glass transition. This enthalpy relaxation (or appearance of the “endothermic hysteresis peak”) is based on the time dependence of the glass transition: if the melt is cooled slowly and then reheated fast, an endothermic hysteresis peak appears on the high-temperature side of the glass transition (Fig. 3.23 for polystyrene). The majority of artificial fibers are semicrystalline. Menczel and Wunderlich (1981) noticed with conventional DSC measurements that for semicrystalline polymers the endothermic hysteresis peak is missing when the sample is cooled much slower than reheated later. Later Menczel (2020b) found that the hysteresis peak is not completely missing and can be seen on the nonreversing heat flow signal when the sample is reheated in modulated mode: the semicrystalline sample is cooled slowly from above the glass transition and reheated at a higher rate in modulated mode (up to a heating rate of 10 C/min). Such DSC curves (the nonreversing signal) are shown in Fig. 3.24. The magnitude of the hysteresis peaks is smaller in the semicrystalline fiber due to the presence of the rigid amorphous phase and the smaller amorphous fraction. The hysteresis peak is more intense for the 1 C/min cooling rate than for the 10 C/min cooling rate, as expected. The fact that the hysteresis peak can be found even for the semicrystalline fibers can be used to identify the

50

Thermal Analysis of Textiles and Fibers

Figure 3.24 The nonreversing heat flow curves indicating the endothermic hysteresis peaks for amorphous PET and semicrystalline drawn PET fiber. The amorphous sample was quenched from 300 C by immersing the pan containing the sample into liquid nitrogen; then the quenched amorphous sample was heated beyond the glass transition temperature to a temperature lower than the starting temperature of cold crystallization (to B100 C) and cooled from this temperature at a rate of 1 and 10 C/min. The semicrystalline sample was crystallized by cooling the sample at a rate of 1 C/min to room temperature. Then this sample was heated to 100 C, and cooled at 1 and 10 C/min to room temperature. Finally, all samples were heated in modulated mode (amplitude/period of 6 1.0 C/40 s) at an underlying heating rate of 10 C/min. The endothermic hysteresis peaks are well visible in the nonreversing heat flow signal. The samples cooled at a rate of 1 C/min display an intensive hysteresis peak because of the slow cooling rate. The samples cooled at a rate of 10 C/min display a hysteresis peak because the samples are being annealed during the heating. The hysteresis peaks are smaller in the semicrystalline drawn fiber due to the smaller amount of the mobile amorphous fraction and the presence of the rigid amorphous fraction. TA Instruments Q2000 DSC (Menczel, 2020b).

glass transition, where it is extremely difficult to see the heat capacity increase. When one takes a look at Fig. 3.25A and B, the glass transition can be easily seen for the amorphous (as-spun) fiber (measurements with nonmodulated, conventional DSC). When the sample is semicrystalline but not oriented (i.e., bulk crystallized without orientation, same figures), the glass transition sometimes is still well visible, although often blurred and broadened. The magnitude of the heat capacity jump for semicrystalline samples is smaller due to the smaller amount of the mobile amorphous fraction. ΔCp can be used to estimate the amount of the mobile amorphous fraction. For PET we know that the ΔCp of the amorphous sample is B78 J/ (mol  C) (Menczel and Wunderlich, 1981). Also, ΔCp is proportional to the amount of the mobile amorphous fraction. The sum of the mobile amorphous fraction calculated from ΔCp and the crystallinity calculated from the heat of fusion is almost never 100%. The difference between 100% and the crystalinity 1 the mobile

Figure 3.25 The DSC curves indicating the glass transition in samples of different physical forms. (A) amorphous and semicrystalline glass transitions poly(ethylene naphthalate) (PEN) (Menczel, unpublished results); (B) DSC curves of amorphous and semicrystalline samples of PET (Menczel, unpublished results); (C) drawn PET fiber, first and second heating: the glass transition is around 80 C, melting point is approximately 270 C (recorded in free-toshrink state): the glass transition is hardly visible in the first heating because of the presence of the oriented amorphous fraction but it can be noted because shrinkage (exothermic peak) starts at the glass transition. The glass transition is better visible in the second heating. TA Instruments Q2000 DSC (Menczel, 2020c). Source: (A) From Menczel, J.D., Prime, R.B., 2009. Thermal Analysis of Polymers, Fundamentals and Applications. John Wiley and Sons. Reprinted with permission from Wiley.

52

Thermal Analysis of Textiles and Fibers

amorphous fraction determines the rigid amorphous fraction: this is mostly the crystal
amorphous interface (Menczel and Wunderlich, 1981; Menczel, 2020a,b,c). As mentioned earlier, it is difficult to see the glass transition on the DSC curves of drawn fibers. The reason for this is the presence of the oriented amorphous fraction. The oriented amorphous fraction exhibits unfreezing in a very broad temperature range, so it is not easily visible. But the abovementioned enthalpy relaxation phenomenon can help to find the glass transition as shown in Fig. 3.26 for PP fiber. For easier determination of the glass transition, one needs to turn to DMA and to use the temperature dependence of tan δ or Ev, the loss modulus (see Chapter 6: Dynamic mechanical analysis in fiber research). Another recently recognized convenient method for measuring the glass transition of drawn fibers is modulated temperature thermomechanical analysis (MT-TMA) (see Fig. 5.3). Also, it is easier to see the glass transition in constrained DSC measurements than in free-to-shrink measurements. The glass transition in free-to-shrink DSC measurements often can be “suspected” only by the presence of the shrinkage exotherm. Modulated temperature DSC experiments can help to find the glass transition also when it is overlapped by some other transitions. Fig. 3.27 shows the MT-DSC curves for drawn Nylon 66 fiber. As can be seen on the total heat flow signal, water evaporation overlaps the glass transition, but the reversing heat flow signal clearly shows the glass transition at around 45 C (Tg is much lower than for dry Nylon 66 because water plasticizes it). It is also noteworthy to mention that the glass transition temperature is somewhat higher in constrained measurements than in free-toshrink measurements. Similarly, when the fiber sample is melted in constrained mode (first heating) and then crystallized and reheated (second heating), Tg in the

Figure 3.26 The endothermic hysteresis peaks of polypropylene fiber cooled from 75 C at various rates (the cooling rates (CR) are indicated at each curve). After cooling, the samples were heated in modulated mode ( 6 1.0 C/40 s) at an underlying heating rate of 10 C/min. Similar hysteresis peaks are observed in the nonreversing heat flow curves as in Fig. 3.24 for PET. TA Instruments Q2000 DSC (Menczel, 2020b).

Differential scanning calorimetry (DSC) in fiber research

53

Figure 3.27 Heating MT-DSC curves of Nylon 66 fiber. The fiber contained a large amount of water, and evaporation of water overlaps the heat capacity jump of the glass transition in the total heat flow signal. The water evaporation is in the nonreversing heat flow signal, but the heat capacity jump of the glass transition can be well seen in the reversing heat flow signal at approximately 45 C. TA Instruments Q2000 DSC (Menczel, 2020c).

Figure 3.28 The glass transition of polypropylene fiber when the measurement is carried out in the constrained (fixed length) mode. The sample was melted, then crystallized at a cooling rate of 10 C/min, and reheated at a rate of 10 C/min. The glass transition during the first heating (sample with higher orientation) is higher than during the second heating (orientation effect). TA Instruments Q2000 DSC (Menczel, 2020c).

first heating is higher than in the second heating obviously due to the higher orientation of the fiber during the first heating (see Fig. 3.28 for PP fiber): the melting during the first heating introduces some disorientation in the fiber. Fig. 3.29 shows the various DSC curves of as-spun PET fiber (first heating
cooling
second heating). The upper curve in Fig. 3.29A is the so-called first

(Continued)

Differential scanning calorimetry (DSC) in fiber research

55

L

heating, this is the DSC curve of the as-received as-spun fiber. There is a heat capacity jump at around 80 C, this is the glass transition of PET. At somewhat higher temperature, an exothermic peak can be seen, this is the cold crystallization indicating that the as-received sample was largely amorphous. If a cold crystallization peak is seen on a DSC heating curve of a polymer, this indicates that the sample was cooled very fast from the melt (quenched), therefore it could not crystallize due to kinetic reasons. On cooling, it passed through the glass transition thus remained largely amorphous. The macromolecular segments of such a sample are frozen below the glass transition (i.e., vibrational, and perhaps, some rotational motion is permitted only). Above the glass transition temperature, the translational motion of the macromolecular segments starts up, and the sample starts crystallizing. Then, at even higher temperatures the just-crystallized portions melt as can be seen from a melting peak at around 260 C
270 C. After the first heating the sample is cooled at a rate of 5 C/min, and the obtained DSC curve is shown in Fig. 3.29B. This curve has an exothermic peak, which corresponds to the melt crystallization of PET. Finally, when this sample is reheated (second heating), we can see a glass transition (Tg) due to the unfreezing of the mobile amorphous fraction and later a melting peak (Fig. 3.29C). There is no cold crystallization, because the sample was cooled relatively slowly. These DSC curves indicate how “fingerprinting” can be used from DSC curves. Cold crystallization indicates that the sample was cooled at high rate (quenched). There are several polymers showing cold crystallization after high cooling rate, which are PET; poly(ethylene naphthalene) (Saw et al., 1997); Nylon M5T [poly(2-methylpentamethylene terephthalamide)] (Menczel et al., 1996, Menczel, 2020a); poly(trimethylene terephthalate) (Brown et al., 1998). On the other hand, other frequently used polymers, such as PE, PP, Nylon 6, polyacrylonitrile do not show cold crystallization, at least at cooling rate

Figure 3.29 First heating, cooling, and second heating of an as-spun fiber (PET): (A) first heating: heating DSC curve of the as-received (amorphous) sample. The first transition is the glass transition, the glass transition temperature is Tg, the heat capacity jump at the glass transition is ΔCp [J/(g  C) or J/(mol  C)]; the exothermic peak indicates cold crystallization: Tcco is the starting temperature of cold crystallization, Tccp is the peak temperature of cold crystallization (in the present case there are two overlapping cold crystallization peaks: they correspond to two nucleation mechanisms, heterogeneous and homogeneous); the third transition is the melting: Tmp is the peak temperature of melting (the temperature at the maximum rate of the melting), and Tm is the melting point. The heat of cold crystallization cannot be determined to a high degree of accuracy because the cold crystallization peak is overlapped by the melting peak (some cold crystallization still may be going on, when the melting of the most imperfect crystallites starts). Therefore the two peaks (cold crystallization and melting) can be integrated with one baseline only, the total heat thus obtained is more-or-less proportional to the crystallinity of the as-received sample. (B) Cooling: melt crystallization curve during cooling; Tco is the starting temperature of melt crystallization, Tcp is the peak temperature of melt crystallization. (C) Second heating: reheating of the melt crystallized sample, the designations are the same as in Fig. 3.29A. TA Instruments Q2000 DSC (Menczel, unpublished results).

56

Thermal Analysis of Textiles and Fibers

available in commercial DSCs. At the same time, we need to mention, that all semicrystalline polymers are likely to exhibit cold crystallization when cooled fast enough like in the fast-scan DSC of Mettler Toledo (Flash 1 DSC), which is in the order of 200,000 C/min. This simply indicates that all semicrystalline polymers can be quenched to an amorphous state at sufficiently high cooling rates. The heating rate is an important parameter in DSC experiments. The results may be totally different with different heating rates. If one heats the sample slowly, the conditions will be right for recrystallization or reorganization during melting. If at slow heating one obtains a double-melting peak, there is a good chance that this was caused by reorganization during melting. It is very easy to decide whether the double-melting peak originates from reorganization during melting: the sample needs to be heated at two different heating rates. If the ratio of the amplitudes of the higher temperature peak to the lower temperature peak increases with decreasing heating rate, the peak doubling is caused by recrystallization (expressions “reorganization during melting” and “recrystallization” are almost synonyms). When the sample is heated fast, time for recrystallization is not sufficient, so with increasing heating rate the higher temperature melting peak (that is the melting of the reorganized crystallites) continuously decreases. At the same time, superheating phenomena may step in (increase of Tm with increasing heating rate). But it is important to remember that superheating is virtually nonexistent for lamellar or fringed micellar crystals. Only extended-chain crystals or highly oriented crystals at fixed length (constrained measurements) can superheat. Superheating usually appears at heating rates higher than 10 C/min, but Illers insisted that superheating may exist in drawn fibers at heating rates just over 0.5 C/min. A detailed study of superheatingreorganization for tire cord PET fiber was carried out by Miyagi and Wunderlich (1972) as a function of sample preparation conditions. Jaffe (1978) carried out heating rate
dependent DSC measurements on unrestrained as-spun PP fibers. According to his observations, the reorganization during melting showed considerable dependence on the spinning conditions. The author observed higher crystallization rates in high-stress-spun fibers because of the nucleating effect of the oriented regions (Fig. 3.30). DSC melting experiments are useful to determine the crystallinity of fibers. The crystallinity is usually (but not accurately) determined by measuring the heat of fusion of the fiber and dividing it by the so-called equilibrium heat of fusion (ΔHf ) of the given polymer. The ΔHf is usually taken from various databases, such as the ATHAS Databank created by Wunderlich (https://www.google.com/ #q 5 Springer%2C 1 heat 1 capacity 1 of 1 polymers) or Wunderlich (1980). However, the equilibrium heat of fusion reported for polymers reflects the heat of fusion of the 100% crystalline polymer at the equilibrium melting point. So, the determination described earlier can cause some error, because the heat of fusion of the equilibrium (extended-chain) crystals depends on temperature. Therefore the readers are referred to the software developed by Vincent Mathot and sold by the Perkin-Elmer Corporation. We also need to mention that the DSC determined crystallinity values are almost always different than the crystallinities determined by X-ray diffraction (XRD). XRD usually ignores the crystallites smaller than a certain

Differential scanning calorimetry (DSC) in fiber research

57

Figure 3.30 Time necessary to reach the maximum crystallization rate (tmax) as a function of melt temperature for polypropylene yarns spun under different stresses: A, fiber spun under high stress; B, fiber spun under low stress. Perkin-Elmer DSC-1B instrument; isothermal crystallization at 130 C, melt holding time 5 min. Source: From Jaffe, M., 1978. Thermal Methods in Polymer Analysis. S.W. Shalaby (Ed.). Franklin Inst. Press, Philadelphia, PA. Reprinted with permission from Franklin Institute Press.

critical value (usually 5 nm), while DSC takes into account the melting of all the crystallites. Menczel (2020c) studied in detail how MT-DSC and MT-TMA measurements can be applied to study the behavior of fibers in constrained and free-to-shrink measurements for PP, Nylon 66, polyacyylonitrile and PET fibers. He found that the nonreversing heat flow gives good indication about the presence of reorganization: for drawn fibers he always found reorganization during melting in free-to-shrink measurements and absence of reorganization in constrained measurements. He observed that not only the melting point but also the glass transition temperature is always higher in constrained measurements of drawn fibers. He also found that the glass transition is much better visible in MT-TMA measurements then by DSC. Constrained and free-to-shrink melting and crystallization experiments indicated that Nylon 66 fibers retain most of the orientation of the polymeric chains when melted once in constrained state: once this fiber melted and crystallized, its melting point on the second heating is not lower than on the first heating, but the melting peak broadens (Fig. 3.31A). Also, the crystallization temperature of the constrained sample is significantly higher than that of the free-to-shrink sample indicating preserved orientation above the melting point (Fig. 3.31B). PP shows totally different behavior. Its melting point decreases significantly on the second heating (Fig. 3.32A), and the crystallization temperature recorded for the constrained and free-to-shrink samples are identical (Fig. 3.32B). This indicates that the chain orientation is almost totally lost on melting. PET exhibits some “medium” behavior: its melting point on the second heating decreases considerably compared to the first heating, but the crystallization temperature of the constrained sample is somewhat

58

Thermal Analysis of Textiles and Fibers

(A) –0.0 1st heating

–0.4 –0.6 –0.8 –1.0

0

(B)

2nd heating

Endotherm

Heat flow (W/g)

–0.2

50

100 200 150 Temperature (°C)

300

2.0

Free-to-shrink

1.5

1.0

Exotherm

Heat flow (W/g)

250

Constrained

0.5

1.0 50

100

150 Temperature (°C)

200

250

Figure 3.31 Melting and crystallization of drawn Nylon 66 fiber. (A) The melting curves in constrained state: first heating: as-received fiber, second heating: melting after crystallizing at a rate of 10 C/min. Heating rate 5 10 C/min. (B) The crystallization in constrained and free-to-shrink measurements (solid line: constrained; dashed line: free-to-shrink). The sample was heated to 290 C in fixed-length (constrained) mode, cooled down at 10 C/min cooling rate, reheated to 290 C at a rate of 10 C/min. TA Instruments Q2000 DSC (Menczel, 2020c).

higher than that of the free-to-shrink sample. This means that some chain orientation is preserved on melting. So, in the constrained measurements, it was found that the loss of orientation during constrained melting is a continuous process, for some fibers it is not destroyed: quasinematic state is formed (partial preservation of the oriented regions). The orientation is preserved in drawn nylon fibers to the highest degree, and this is likely to be caused by hydrogen bonding in Nylon. For several polymers, especially for PET, cold crystallization can be a good fingerprinting technique. Fig. 3.33 shows the cold crystallization of PET for various

Differential scanning calorimetry (DSC) in fiber research

59

Figure 3.32 Melting and crystallization of drawn polypropylene fiber. (A) The melting curves in constrained state: first heating: as-received fiber (solid curve), second heating: melting after crystallizing at a rate of 10 C/min (dashed curve). Heating rate 5 10 C/min. (B) The crystallization in constrained and free-to-shrink measurements (solid line: constrained; dashed line: free-to-shrink). The sample was heated to 230 C in fixed-length mode, cooled down at 10 C/min cooling rate, reheated to 230 C at a rate of 10 C/min. TA Instruments Q2000 DSC (Menczel, 2020c).

spinning conditions (Heuvel and Huisman, 1978). It can be seen that increasing windup speed of the PET yarn enhances the segmental orientation in the as-spun yarn in the melt, and the increased orientation behaves as a nucleating agent, it speeds up the crystallization process (see also Smith and Stewart, 1974). The peak temperature of cold crystallization (Fig. 3.33) decreases linearly with increasing orientation, so it can be used as a parameter characterizing the orientation. Similar measurements were performed by Jaffe (1978) on PP fibers produced with different melt stresses at different melt temperature and melt time in the DSC. Jaffe

60

Thermal Analysis of Textiles and Fibers

Figure 3.33 DSC heating curves of as-spun poly(ethylene terephthalate) yarns from different windup speeds. The yarns were spun to constant count of dTex 167/30. The spin temperature was 290 C, 250 μm hole size. DuPont (TA Instruments) 990/910 DSC, heating rate 20 C/min. Source: From Heuvel, M., Huisman, R., 1978. J. Appl. Polym. Sci. 22, 2219. Reprinted with permission from Wiley.

systematically changed the melt temperature and the time in the melt and carried out isothermal crystallization at 130 C (see Fig. 3.30). He observed that the crystallization was faster in high-stress-spun PP fibers obviously indicating the nucleating effect of the oriented regions. These results underline the importance of knowing and adjusting the melt history during production. Heat setting is an important step in fiber manufacturing. Berndt et al. (1973) and Berndt and Bossmann (1976) performed a detailed study on clarifying the nature of this endothermic peak in commercial PET fiber. The intensity of this peak depends on the annealing temperature, time, tension, and media affecting the yarns in aftertreatments such as heat setting or dyeing. The position of this endothermic peak increases with increasing dwell time and increasing heat-setting temperature. By now, we know that annealing at any temperature (and heat setting is an annealing process) below the melting peak (but above the glass transition temperature)

Differential scanning calorimetry (DSC) in fiber research

61

Figure 3.34 DSC curves of flame-retarded poly(ethylene terephthalate) yarns in different atmospheres: N2, solid line; air, dashed line. DuPont (TA Instruments) 990/910 DSC, sample mass 15 mg, heating rate 10 C/min. Source: From Hassel, R.L., 1977. Am. Lab. (Fairfield, Conn.) 9, 35. Reprinted with permission from International Scientific Communications, Inc.

produces a small endothermic peak that reflects melting of small defective crystals crystallized slowly during the annealing process. We also know that this phenomenon is not related to fibers alone, but it is a general phenomenon in DSC measurements of semicrystalline polymers. DSC sometimes is used to evaluate the thermal stability of polymers, including fibers. Decomposition can be exothermic or endothermic depending on the atmosphere. Decomposition in inert atmosphere (such as nitrogen), which is the most frequently used purge gas in DSC, is often endothermic, while it is exothermic in air. Hassel (1977) studied flame-retarded PET yarns heated to high temperature in air and nitrogen (see Fig. 3.34). Of course, DSC is a preferred method for determining the thermal history of any polymer, although TMA sometimes can be more sensitivie. The reason of this is simple: the processing conditions are far from equilibrium, therefore the obtained thermal properties will reflect the thermal and mechanical history, including drawing of the fibers. To illustrate this for fiber drawing, Fig. 3.35 shows the DSC curves of as-spun (original) and drawn PEEK fibers. The as-spun fiber is largely amorphous, because of the fast cooling. Therefore it has an intense glass transition followed by cold crystallization, and finally melting. In the drawn fibers the glass transition is hardly visible (this is the consequence of the presence of the oriented amorphous phase), and there is no cold crystallization, because the fibers are partially crystalline. Also, DSC is a preferred method to determine the crystallinity of the fibers. XRD may give somewhat different results than DSC, because XRD may not feel the presence of smallest crystallites in polymers. Fig. 3.36 shows the

62

Thermal Analysis of Textiles and Fibers

Figure 3.35 DSC curves of as-spun and different drawn PEEK fibers. The draw ratios are indicated at each curve. DuPont (TA Instruments) 990/910 DSC, heating rate is 20 C/min. Source: From Song, S.S., White, J.L., Cakmak, M., 1989. Sen’i Gakkaishi 45 (6), 243. Reprinted with permission from Sen’i Gakkaishi.

crystallinity of PEEK [poly(ether ether ketone)] fibers drawn at different temperatures determined by integration of the DSC curves. Finally, some words on the use of DSC for studying the behavior of natural fibers (Chapter 7 in this book describes thermal analysis of natural fibers). The most important natural fibers are wool, silk, and cotton. As it is well known, one of the major distinguishing properties of natural fibers is their ability to pick up large amounts of water. Therefore it is natural that one of the major applications of DSC is to characterize the state of water in these fibers. The two possible states of water in a polymer are bound water and free water (there may be several types of bound water, the two most often met are tightly bound water and loosely bound water). Fig. 3.37 shows that DSC can differentiate between free water and bound water in Merino wool, cotton, and silk. Basch and Lewin (1975) carried out DSC measurements with the purpose of clarifying how the physical structure and the presence of various chemical agents affect the decomposition of cellulosics (see Fig. 3.38). They treated cellulosic fabrics with sulfuric acid, ammonium sulfamate, ammonium bisulfate, urea
phosphoric acid, and diammonium phosphate and recorded the DSC curves of these systems: the peak temperature of decomposition varied depending on the additive. Despite the absence of reversible thermal transitions (except the glass transition), DSC can be very helpful for characterizing natural fibers. Chemically untreated silk has the glass transition at 160 C as evidenced by DMA measurements. Tsukada (1988) and Tsukada et al. (1993) studied grafting silk fibers with 2-hydroxyethyl methacrylate (HEMA, Fig. 3.39). The authors proved that the HEMA treatment leads to improved thermal stability of silk. Fig. 3.39 indicates that untreated silk has one endothermic transition at approximately 325 C corresponding

Figure 3.36 Crystallinity of PEEK fibers for various draw ratios determined by integration of the DSC curves. The drawing temperatures were 160 C, 200 C, and 230 C. DuPont (TA Instruments) 990/910 DSC, heating rate is 20 C/min. Source: From Song, S.S., White, J.L., Cakmak, M., 1989. Sen’i Gakkaishi 45 (6), 243. Reprinted with permission from Sen’i Gakkaishi.

Figure 3.37 Cooling and heating DSC curves showing the freezing of water sorbed on Merino wool, cotton yarn, and silk and the melting of ice at different levels of water content. The runs were made on a Rigaku-Denki DSC-10A. The samples were ground to fine powder to eliminate the effects of orientation, then the powder was dried in vacuum at 20 C for 2 weeks. Heating and cooling rates were 5 C/min. Source: From Sakabe, H., Ito, H., Miyamoto, T., Inagaki, H., 1987. Text. Res. J. 57 (1), 66. Reprinted with permission from Textiles Research Institute.

64

Thermal Analysis of Textiles and Fibers

Figure 3.38 Effect of flame-retardant systems and fine structure on the DSC curves of cellulosic fabrics. (A) Cotton fabric sulfated with sulfuric acid; (B) rayon fibers sulfated with ammonium sulfamate; (C) ramie impregnated with ammonium bisulfate; (D) cotton fabric phosphorylated with urea
phosphoric acid; (E) ramie impregnated with diammonium phosphate; (F) rayon phosphylated with urea
phosphoric acid. DuPont (TA Instruments) 900 DSC, heating rate 10 C/min. Source: From Basch, A., Lewin, M., 1975. Text. Res. J. 45, 246. Reprinted with permission from Textile Research Institute.

to thermal decomposition of silk fibroin with oriented β0 -configuration. The other two peaks (at 280 C and 425 C) on Fig. 3.39 represent the thermal decomposition of methacrylamide (MAA) and HEMA. Wool has one reversible thermal transition at 160 C
180 C correlated with the onset of side-chain motion. This is a “pseudo-glass transition.” There is an endothermic peak ascribed to some chemical changes at 171 C
194 C. The major transition is a double endothermic transition at 235 C and 243 C: at this temperature range wool loses the α-helical content and shrinks. Finally, thermal decomposition takes place at 292 C
321 C. Konda et al. (1973) studied the 230 C endothermic

Figure 3.39 DSC curves of untreated and grafted silk fibers. (A) untreated silk fiber; (B) HEMA/MAA mixture of 0/100 with graft yield of 59%; (C) HEMA/MAA mixture of 20/80 with graft yield of 61%; (D) HEMA/MAA mixture of 0/100 with graft yield of 120%. Grafting was done at 75 C
80 C using potassium persulfate as initiator. The curves were recorded on a Rigaku-Denki DSC-10A. Source: From Tsukada, M., 1988. J. Appl. Polym. Sci. 35, 2133. Reprinted with permission from Wiley.

Figure 3.40 DSC curves of various drawn and relaxed wool keratin fibers. The drawing was carried out in water at 20 C, the relaxation performed in cuprammonium solution. Measurements were done by a Rigaku-Denki DSC 10A; sample mass was B3 mg, purge gas N2. Source: From Konda, A., Tsukada, M., Kuroda, S., 1973. J. Polym. Sci., Polym. Lett. Ed. 11, 247. Reprinted with permission from Wiley.

66

Thermal Analysis of Textiles and Fibers

Figure 3.41 DSC and TGA curves of untreated and flame-retarded cotton: (A) untreated cotton; (B) cotton pyrovatex; (C) Durelle. The TGA curves were recorded on a Stanton Model HT-SM thermobalance. The sample mass was 100
200 mg, purge gas: air, heating rate 6 C/min. Source: From Bingham, M.A., Hill, B.J., 1975. J. Therm. Anal. 1, 347. Reprinted with permission from Springer.

event for variously drawn wool fibers (Fig. 3.40). With the help of XRD, the peak was assigned to disappearance of the α-form of wool keratin. Of course, DSC can be used and is being used to characterize thermal stability of various natural fibers especially for characterizing polymer flammability. In such cases, TGA is also a very helpful technique. Fig. 3.41 shows the DSC and TGA curves of untreated and flame-retarded cotton (Bingham and Hill, 1975) for

Differential scanning calorimetry (DSC) in fiber research

67

evaluating thermo-oxidative decomposition. There are major differences in both DSC and TGA curves under thermo-oxidative degradation conditions. Interpretation of thermal analysis recordings is not always easy, and often DSC recording needs to be supplemented with results from other thermal analysis and analytical techniques, such as XRD. We mentioned that it is very difficult to determine the glass transition of drawn fibers by DSC, because the heat capacity jump is very shallow and broad, and often even experienced eyes cannot notice this transition. In such cases, DMA or MT-TMA or can be helpful (see Fig. 5.3). Similarly, when evaluating the thermal or thermo-oxidative degradation of fibers or other polymers, TGA measurements can be helpful for interpretation of the DSC recordings. In this chapter, we described the basics and application of DSC to fiber research. It was presented that DSC is one of the most versatile thermal analysis techniques, which can be applied to any fiber for studying its thermal properties. DSC gives very useful data that can be used to deduct structural information of fibers.

References Ancsin, J., 1985. Metrologia 21 (1), 7. Archer, D.G., Rudtsch, S., 2003. J. Chem. Eng. Data 48 (5), 1157. ASTM Standard E2069-00, 2012. Standard Test Method for Temperature Calibration on Cooling of Differential Scanning Calorimeters. Aston, J.G., Szasz, G.J., Finke, H.L., 1943. J. Am. Chem. Soc. 65, 1135. ATHAS Databank. ,https://www.google.com/#q 5 Springer%2C 1 heat 1 capacity 1 of 1 polymers.. Bares, V., Wunderlich, B., 1973. J. Polym. Sci., Polym. Phys. Ed. 11, 861. Basch, A., Lewin, M., 1975. Text. Res. J 45, 246. Berndt, H.-J., Bossmann, A., 1976. Polymer 17, 241. Berndt, H.-J., Schultz, H., Heidemann, G., 1973. Melliand. Textilber. 54, 773. Bingham, M.A., Hill, B.J., 1975. J. Therm. Anal. 1, 347. Boller, A., Jin, Y., Wunderlich, B., 1994. J. Therm. Anal. 42, 307. Brown, H., Hwo, C., Grebowicz, J., 1998. Proc. 8th TANDEC Meeting, Knoxville, TN, 1998. Cottrell, T.L., 1954. The Strengths of Chemical Bonds. Butterworth, London. Coursey, J.S., Schwab, D.J., Tsai, J.J., Dragoset, R.A., 2014. Atomic Weights and Isotopic Compositions (Version 3.0), 2010. National Institute of Standards and Technology, Gaithersburg, MD. Emsley, J., 2011. Nature’s Building Blocks: An A-Z Guide to the Elements, second ed. Oxford University Press, New York. Gagniuc, P.A., 2017. Markov Chains: From Theory to Implementation and Experimentation. John Wiley & Sons, NJ, ISBN: 978-1-119-38755-8pp. 46
59. Gaur, U., Wunderlich, B., 1980. Macromolecules 13, 445. Gaur, U., Wunderlich, B., 1981. J. Phys. Chem. Ref. Data 10, 1051. Gaur, U., Wunderlich, B., 1982. Computer Applications in Applied Polymer Science, Chapter 21. In: ACS Symposium Series, 197, pp. 355
366. Available from https://doi. org/10.1021/bk-1982-0197. ch021. ISBN13: 9780841207332eISBN: 9780841209305. Hassel, R.L., 1977. Am. Lab. (Fairfield, Conn.) 9, 35.

68

Thermal Analysis of Textiles and Fibers

Haynes, W.M. (Ed.), 2014. CRC Handbook of Chemistry and Physics. 95th ed. CRC Press/ Taylor and Francis, Boca Raton, FL. Heuvel, M., Huisman, R., 1978. J. Appl. Polym. Sci. 22, 2219. Illers, K.-H., 1970. Angew. Makromol. Chem. 12, 89. Ishikiriyama, K., Wunderlich, B., 1997. J. Therm. Anal. 50, 337. Jaffe, M., 1978. In: Shalaby, S.W. (Ed.), Thermal Methods in Polymer Analysis. Franklin Inst. Press, Philadelphia, PA, p. 93. Jaffe, M., Wunderlich, B., 1967. Kolloid.-Z. Z. Polym. 216
217, 203. Jin, Y., Boller, A., Wunderlich, B., 1993. Proc. 22nd North American Thermal Analysis Soc. Conf., Denver, CO, September 19
22, 1993, p. 59. Kaye & Laby Online, 1995. Tables of Physical & Chemical Constants, 16th ed. Version 1.0 (2005) (accessed December 2014). Konda, A., Tsukada, M., Kuroda, S., 1973. J. Polym. Sci., Polym. Lett. Ed. 11, 247. Kurnakov, N.S., 1904. Z. Anorg. Chem. 42, 184. Le Chatelier, H., 1887. Z. Phys. Chem. 1, 396. Mead, W.T., Porter, R.S., 1976. J. Appl. Phys. 47 (C10), 4278. Menczel, J.D., 1994. Thermal Analysis News, Worldwide Perkin-Elmer Thermal Analysis Customer Newsletter (Sept.), pp. 6
7. Menczel, J.D., 1997. J. Therm. Anal. 49, 193. Menczel, J.D., 2020a, to be published in JTAC. Menczel, J.D., 2020b, to be published. Menczel, J.D., 2020c, to be published. Menczel, J.D., Leslie, T.M., 1990. Thermochim. Acta 166, 309. Menczel, J.D., Leslie, T.M., 1993. J. Thermal Anal. 40, 957. Menczel, J.D., Jaffe, M., 2007. Proc. 34th North American Thermal Analysis Soc. Conf. J. Therm. Anal. Calorim. 89 (2), 357. Menczel, J.D., Prime, R.B., 2009. Thermal Analysis of Polymers, Fundamentals and Applications. John Wiley and Sons. Menczel, J.D., Grebowicz, J., 2020, 2021, 2023. Handbook of Differential Scanning Calorimetry, vols. 1, 2 and 3. Elsevier (in preparation). Menczel, J.D., Jaffe, M., Saw, C.K., Bruno, T.P., 1996. J. Therm. Anal. 49, 201. Menczel, J.D., Wunderlich, B., 1981. J. Polym. Sci., Polym. Lett. Ed. 19, 261. Menczel, J.D., Wunderlich, B., 1986. Polym. Prepr., J. Am. Chem. Soc., Div. Polym. Chem. 255. Miller, B., 1971. Thermochim. Acta 2, 225. Miyagi, A., Wunderlich, B., 1972. J. Polym. Sci., Polym. Phys. Ed. 10, 1401. Oswald, H.J., Turi, E.A., Harget, P.J., Khanna, Y.P., 1977. J. Macromol. Sci., Phys. B13 (2), 231. Reading, M., Hourston, D.J., 2006. Modulated Temperature Differential Scanning Calorimetry. Springer, Dordrecht. Roberts-Austen, W.C., 1899. Metallographist 2, 186. Sakabe, H., Ito, H., Miyamoto, T., Inagaki, H., 1987. Text. Res. J. 57 (1), 66. Samuels, R.J., 1975a. J. Polym. Sci., Polym. Phys. Ed. 13, 1417. Samuels, R.J., 1975b. Appl. Polym. Symp. 27, 205. Saw, C.K., Menczel, J., Choe, E.W., Hughes, O.R., 1997. SPE ANTEC ’97, April 27
May 2, 1997, Toronto, vol. II, Materials, p. 916. Schick, C., Mathot, V. (Eds.), 2016. Fast Scan Calorimetry. Springer. Schumm, R.H., Wagman, D.D., Bailey, S., Evans, W.H., Parker, V.B., 1973. Nat. Bureau St. (USA) Technical Notes 270-1 to 270-8. Smith, F.S., Stewart, R.D., 1974. Polymer 15, 283. Smook, J., Pennings, J., 1984. Colloid Polym. Sci. 262, 712.

Differential scanning calorimetry (DSC) in fiber research

69

Song, S.S., White, J.L., Cakmak, M., 1989. Sen’i Gakkaishi 45 (6), 243. Tashiro, K., Naki, Y., Kobayashi, M., Tadokoro, H., 1980. Macromolecules 13, 137. Thomas, L.C., 2006. Modulated DSC Technology. TA Instruments (Publisher). Todoki, M., 1985. Thermochim. Acta 93, 147. Todoki, M., Kawaguchi, T., 1977a. J. Polym. Sci., Polym. Phys. Ed. 15, 1067. Todoki, M., Kawaguchi, T., 1977b. J. Polym. Sci., Polym. Phys. Ed. 15, 1507. Tsukada, M., 1988. J. Appl. Polym. Sci. 35, 2133. Turi, E.A., 1997. Thermal Characterization of Polymeric Materials. Academic Press. Valk, G., 1972. Lenzinger Ber. 33, 1. Valk, G., Berndt, H.J., Heidemann, G., 1971. Chemiefasern 5, 1. Varga, J., 1983. Private communication to J. Menczel. Varga, J., Menczel, J.D., Solti, A., 1979. J. Therm. Anal. 17, 333. Varma-Nair, M., Wunderlich, B., 1990. J. Phys. Chem. Ref. Data 20, 349. Wiesner, E., 1974. Chem. Vlakna 24 (1), 1. Wiesner, E., 1976a. Termanal’76, Celostatna Konf. Term. Anal. [Pr.], 7th, 1976 0-69-0-75. Wiesner, E., 1976b. Chem. Vlakna 26 (3-4), 146. Wunderlich, B., 1980. Macromolecular Physics, vol. 3. Academic Press, New York; London; Toronto; Sydney; San Francisco, CA. Wunderlich, B., 1990. Thermal Analysis. Academic Press, Boston, MA. Zachmann, H.G., 1965. Kolloid-Z. Z. Polym. 206, 25. Zachmann, H.G., Spellucci, P., 1966. Kolloid-Z. Z. Polym. 213, 39.

Further reading Barrau, S., Judovits, L., 1999. Proceedings of the 27th NATAS Conference, p. 284. Danley, R.L., 2003. Thermochim. Acta 395, 201. Danley, R.L., 2004. Thermochim. Acta 409, 111. Dimov, K., Georgiev, J., Bechev, C., 1973. Faserforsch. Textiltech. 24 (8), 337. Hemminger, W., Ho¨hne, G., 1984. Calorimetry Fundamentals and Practice. Verlag Chemie. Mathot, V.B.F., Poel, G.V., Pijpers, T.F.J., 2006. Am. Lab. 38, 21. Mettler Toledo Stare System Manual, Chapter 8, 2004. O’Neill, M.J., 1964. Anal. Chem. 36, 1238. TA Instruments, 1993. DSC 2920 Differential Scanning Calorimeter Operator’s Manual. Wunderlich, B., 2006. Thermal Analysis of Polymeric Materials. Optical Disk.

Thermogravimetric analysis of fibers

4

Joseph D. Menczel Thermal Measurements LLC, Fort Worth, TX, United States

Abstract In this chapter the basics of thermogravimetric analysis (TGA) are explained, and the use TGA in research and characterization of fibers is shown. It is explained how to use TGA for measuring the losses of low-molecular-mass materials during temperature increase and how to characterize thermal and thermo-oxidative stability of organic polymer based fibers. The differences between thermal stabilities of bulk polymers and fibers are illustrated. It is demonstrated how to use TGA for determining the effects of crystallinity and molecular orientation on the thermal stability of fibers. Examples are given to describe the effect of flame retardation on thermal stability and various chemical treatments on thermal and thermo-oxidative stability of natural fibers.

“Thermogravimetric analysis (TGA) is defined as a thermal-analysis technique in which the sample mass is measured as a function of temperature while the sample is subjected to a controlled temperature program in a controlled atmosphere” (Earnest, 1988). TGA can provide information about the physical phenomena accompanied by a mass change, such as phase transitions, including vaporization, sublimation, absorption, adsorption, and desorption. Similarly, TGA can provide information about chemical transformations, such as polymer degradation, chemisorption, dehydration, decomposition, and solid gas and melt gas reactions (e.g., oxidation). TGA measures the mass change of a material as a function of increasing temperature or at a constant temperature isothermally as a function of time in a specified atmosphere (nitrogen, air, some other gas, or in vacuum). The temperature of TGA measurements ranges usually from room temperature to 1,000 C (in high temperature TGA up to 1,500 C or even 2,000 C). Sample masses range from 1 to 20 30 mg. Mass change sensitivity is of 0.001 mg. TGA is often used to determine the content of water or other solvents in polymers and the rate of evaporation of these solvents from polymers. Oxidation reactions can be followed by TGA, because the absorption of oxygen appears as a mass increase. TGA can be used to determine the content of inorganic fillers in polymers by burning off the organic polymer. But the major use of TGA is to determine the thermal (nitrogen and helium atmosphere) or thermo-oxidative (air or oxygen) stability of polymers, because thermal or thermo-oxidative degradation of these Thermal Analysis of Textiles and Fibers. DOI: https://doi.org/10.1016/B978-0-08-100572-9.00004-5 © 2020 Elsevier Ltd. All rights reserved.

72

Thermal Analysis of Textiles and Fibers

Figure 4.1 TGA curves of an extruded polypropylene piece and polypropylene fiber. PerkinElmer TGA7 instrument, purge gas is N2, heating rate 10 C/min. Source: Menczel (unpublished results).

materials is accompanied by mass loss. The TGA instrument is often interfaced with other analytical instruments, most often with a Fourier-transform infrared (FT IR) spectrometer or a mass spectrometer. TGA results depend significantly on the specific surface area of the sample, so these results are very sensitive to the thickness of fibers because of the significant surface-to-volume ratio. Usually the degradation rate constant calculated from these TGA curves increases when the specific surface area increases. Therefore these details of the sample preparation are important, and the following details must be included in the reports of TGA analyses: sample mass, sample preparation, heating rate, and purge gas. Fig. 4.1 shows comparison of the TGA curves of bulk and fiber polypropylene. It is clear that the increase of the specific surface area (from bulk to fiber) has a huge effect on the degradation temperature. For TGA measurements of drawn thin fibers, 2 mg sample mass and 10 C/min heating rate can be standard. When reporting the results of TGA measurements, the mass of the sample, the sample preparation conditions (e.g., grinding the sample), heating rate, and the atmosphere should always be included. When results of various measurements are compared, it is preferable to make the runs with similar conditions. TGA measurements are important for determination of the water content and equilibrium water content of textile fibers. The water pickup from human sweat is an essential property of the fibers. The use of hydrophobic fibers is very inconvenient in human clothing. TGA seems to be a very fast and convenient thermal-analysis technique for this application. Other water-determination techniques, such as Karl Fischer titration and oven drying, are lengthy, and Mansrekar (1973) came to the conclusion that TGA measurements provide the best compromise between accuracy and convenience of the moisture determination. In reality, when a textile piece is given, all you have to do is to load it into the TGA pan, close the furnace, and press the Start button. In 30 minutes, you have a

Thermogravimetric analysis of fibers

73

TGA curve that can be evaluated for the moisture content literally in less than a minute. TGA is mostly used to characterize polymeric materials (including fibers) as a function of chemical composition, although physical variables (such as specific surface area) can play a significant role in the response. TGA’s major use in fiber characterization is the following: G

G

G

G

Characterization of the degradation of the fiber in various (inert or reactive) atmospheres Monitoring the loss of moisture, various solvents, other volatile additives, and/or finish content Characterization of thermal stability of the fiber under the influence of certain additives Determination of degradation kinetics of the fiber (which may be different than the degradation kinetics of the basic polymeric materials due to the huge specific surface area)

Fibers are rarely used for the determination of kinetics of thermal degradation. The mass of the fiber that can be loaded into the TGA pan is very small, and this represents a special difficulty with the TGA measurements of fibers. When loading fibers into the TGA pan, the fiber pieces often fly around and leave the TGA pan due to static electricity. So, whenever the general degradation properties of a polymeric substance need to be determined, bulk polymeric materials are used. Fibers are used as samples when the specific surface area becomes an issue. Basch and Lewin (1973a,b, 1974) studied in detail how crystallinity and the molecular orientation affect the vacuum and oxidative pyrolysis of cellulose. They used very well-characterized cellulosic samples. They showed that crystallinity is a major factor in the stability of cellulose. Using samples of different crystallinities, they determined the apparent activation energy of pyrolysis for the amorphous and crystalline regions. Their results indicated that the apparent activation energy of pyrolysis for the amorphous regions was about 125 kJ/mol, and double this value for the crystalline regions, as shown in Fig. 4.2. The orientation had a complex effect on pyrolysis: it slowed down the reaction in air and accelerated it in vacuum. Fig. 4.3 shows isothermal mass losses for 20% and 50% stretched cellulose fibers. It is clear that stretching increases the thermal stability as expected. The slowdown of the reaction in air was explained by a decreased accessibility of oxygen to the cellulosic structure, while the rate increase in vacuum pyrolysis was explained by the increase in proximity of the cross-link forming free radicals. Polymer flammability is an important area in TGA characterization of fibers especially because of the huge specific surface area of thin fibers. The role of flame retardants is critical here. Bingham and Hill (1975) collected a list of flame retardants and used TGA extensively for comparison of thermogravimetric behavior of various cotton samples with and without flame retardants under thermo-oxidative degradation conditions in air (Fig. 4.4). Carroll-Porczinski (1972) developed an interesting TGA method for flammability evaluation. The author constructed an additive TGA curve for polymer blends from the TGA curves of the individual components and compared this curve to the experimental TGA curve of the blend. If the mass loss of the additive curve is

74

Thermal Analysis of Textiles and Fibers

Figure 4.2 Cellulosic yarns: the apparent activation energy of pyrolysis as a function of cellulose crystallinity as measured by TGA using the Horowitz and Metzger method (Horowitz and Metzger, 1963). DuPont (TA Instruments) 950 TGA, sample mass 5 mg, heating rate 10 C/min. The crystallinity was determined by wide angle X-ray scattering (WAXS). Source: From Basch, A., Lewin, M., 1973a. J. Polym. Sci., Polym. Chem. Ed. 11, 3071, Fig. 1. Reprinted with permission of Wiley.

Figure 4.3 Vacuum pyrolysis experiments of rayon yarns: the mass loss versus time at 151 C for two different yarn orientations (degrees of stretch). DuPont (TA Instruments) 950 TGA. Source: From Basch, A., Lewin, M., 1974. J. Polym. Sci., Polym. Chem. Ed. 12, 2053, Fig. 2. Reprinted with permission of Wiley.

Thermogravimetric analysis of fibers

75

Figure 4.4 Thermo-oxidative degradation of untreated and flame-retarded cotton characterized by DSC and TGA. (A) Untreated cotton, (B) cotton Pyrovatex, and (C) Durelle. The curves were recorded on a Stanton Model HT-SM thermobalance with 100 200 mg sample mass in air. The heating rate was 6 C/min. Source: From Bingham, M.A., Hill, B.J., 1975. J. Therm. Anal. 1, 347, Figs. 1 3. Reprinted with permission of Wiley.

greater than the mass loss on the TGA curve of the actual composite, the flammability performance of the polymer blend was considered synergistic and acceptable. Two examples are shown in Fig. 4.5 as a “good” blend (50/50 wool/Nomex) and a “bad” blend (50/50 wool/polypropylene). Gaulin and McDonald (1971a,b, 1972) used TGA to determine the radiation effect on the pyrolysis of PAN yarns in vacuum (Fig. 4.6): the behavior of asreceived, low-comonomer content polyacrylonitrile (PAN) fiber was compared to the 60Co-irradiated PAN, 60Co- and e-irradiated PAN, and 60Co-irradiated and

Figure 4.5 TGA curves of wool/Nomex and wool/polypropylene composite fabrics. The solid curves indicate the actual curves of component fibers and the composite. The dashed curves are the additive curves constructed of the component fiber results. TR-02 STA TGA instrument. Purge gas: air (50 mL/min flow rate), 10 C/min heating rate. Source: From Carroll-Porczynski, C.Z., 1972. Proc. Third ICTA Conference 1971, Davos, vol. 3, p. 273; Figs. 1 3. Reprinted with permission of Birkhauser Verlag.

Figure 4.6 TGA curves recorded in vacuum of polyacrylonitrile yarns after different irradiation treatments. The basic PAN yarn was an American Cyanamid produced lowcomonomer content yarn. (A) As-received yarn; (B) 60Co-irradiated yarn; (C) 60Co and electron irradiated yarn; (D) 60Co-irradiated and pre-oxidized yarn. 5 C/min heating rate. Source: From Gaulin, G.A., McDonald, W.R., 1971a. Air Force Rep. SAMSO-TR-71-61, vol. 1; Gaulin, G.A., McDonald, W.R., 1971b. Space Missile Syst. Organ, Rep. TR-0059 (6250-40)-5, vol. 1, Fig. 2. Public domain document.

Thermogravimetric analysis of fibers

77

Figure 4.7 (A) TGA curves of predried keratin fibers. (a) Merino 70’s (top: semi-processed longer fiber pieces); (b) Lincoln 44’s (top); (c) Blackface 36’s (top); (d) cashmere (noil, short fiber pieces); (e) angora fur. Sample mass in the TGA measurements: 10 mg. The fiber pieces were cut to pass through 60-mesh sieve. During the measurements continuous vacuum, heating rate was 4 C/min. (B) Derivative curves (DTGA) of the curves shown in Fig. 4.11A in Crighton and Hole (1976). Source: From Crighton, J.S., Hole, P.N., 1976. Proc. Int. Wolltextil-Forschungskonf., 5th, 1975, p. 499, Figs. 1 and 2. Reprinted with permission of Wiley.

preoxidized PAN yarns. Also, TGA was very successfully used in the development of graphite-fiber precursors. Fitzer and Mueller (1971, 1973) and Grassie and McGuchan (1970, 1972) used TGA for completing kinetic studies of PAN pyrolysis for graphite-fiber production. Crighton and Hole (1976) applied TGA to characterize thermal degradation of various wool samples (Fig. 4.7). TGA is extensively used for characterization of natural fibers of both animal and plant origin as is shown in Fig. 4.7 by Crighton and Hole (1976) for various wool

78

Thermal Analysis of Textiles and Fibers

Figure 4.8 TGA curves of various untreated and chemically treated cotton fabrics recorded in oxygen: (U) untreated; (S) scoured; (M) mercerized; (A) acetylated; and (B) benzoylated. Perkin-Elmer TGA7 instrument. Source: From Subramanian, V., Vasugi, N., 1989. J. Appl. Polym. Sci. 38, 207, Fig. 1. Reprinted with permission of Wiley.

samples. These curves show that the derivative TGA curves are often more effective for characterization of these samples. Subramanian and Vasugi (1989) used TGA to characterize the effect of chemical treatment on thermal stability of cotton (Fig. 4.8). In general, acetylation and benzoylation improved the various physical properties of cotton; among other things, they lead to increased thermal stability. Day et al. (1987) used thermogravimetry to characterize the effect of flame retardation on thermal stability of PET fabrics. They recorded the TGA curves of the basic poly(ethylene terephthalate) (PET) fabric and PET fabrics containing poly(4bromostyrene), poly(vinyl bromide), and poly(vinylidene bromide) (Fig. 4.9). TGA results were compared to flammability limits, hydrogen bromide evolution, and the limiting oxygen index data in order to estimate the efficiency of flame retardation of the mentioned bromo-containing compounds. All these additives released HBr during the pyrolysis. However, the release of HBr from the fabrics containing poly (vinyl bromide) and poly(vinylidene bromide) took place at temperatures too low for efficient flame retardation.

Thermogravimetric analysis of fibers

79

Figure 4.9 TGA curves of untreated and flame-retarded PET fabrics. Solid curve: untreated PET fabric; dashed curve: PET fabric treated with poly(4-bromostyrene); dashed-dotted curve: PET fabric containing poly(vinyl bromide); dotted curve: PET fabric containing poly (vinylidene bromide). DuPont Instruments (TA Instruments) 1090/951 TGA, heating rate 5 C/min. Source: From Day, M., Suprunchuk, T., Cooney, J.D., Wiles, D.M., 1987. J. Appl. Polym. Sci. 33, 204, Fig. 1. Reprinted with permission of Wiley.

References Basch, A., Lewin, M., 1973a. J. Polym. Sci., Polym. Chem. Ed. 11, 3071. Basch, A., Lewin, M., 1973b. J. Polym. Sci., Polym. Chem. Ed. 11, 3095. Basch, A., Lewin, M., 1974. J. Polym. Sci., Polym. Chem. Ed. 12, 2053. Bingham, M.A., Hill, B.J., 1975. J. Therm. Anal. 1, 347. Carroll-Porczynski, C.Z., 1972. Proc. Third ICTA Conference 1971, Davos, vol. 3, p. 273. Crighton, J.S., Hole, I.N., 1976. Proc. Int. Wolltextil-Forschungskonf., 5th, 1975, p. 499. Day, M., Suprunchuk, T., Cooney, J.D., Wiles, D.M., 1987. J. Appl. Polym. Sci. 33, 204. Earnest, C.M., 1988. Compositional Analysis by Thermogravimetry, ASTM STP 997. ASTM International, Philadelphia, PA. Fitzer, E., Mueller, D.J., 1971. Makromol. Chem. A4, 117. Fitzer, E., Mueller, D.J., 1973. Polym. Prepr., Am. Chem. Soc., Div. Polym. Chem. 14 (1), 386. Gaulin, G.A., McDonald, W.R., 1971a. Air Force Rep. SAMSO-TR-71-61, vol. 1. Gaulin, G.A., McDonald, W.R., 1971b. Space Missile Syst. Organ, Rep. TR-0059(6250-40)5, vol. 1. Grassie, N., McGuchan, R., 1970. Eur. Polym. J. 6, 1277. Grassie, N., McGuchan, R., 1972. Eur. Polym. J. 8, 865. Mansrekar, T.G., 1973. Chem. Color 5, 186. Subramanian, V., Vasugi, N., 1989. J. Appl. Polym. Sci. 38, 207.

Further reading Basch, A., Lewin, M., 1975. Text. Res. J. 45, 246.

Thermomechanical analysis of fibers

5

Joseph D. Menczel1 and Michael Jaffe2 1 Thermal Measurements LLC, Fort Worth, TX, United States, 2New Jersey Innovation Institute, University Heights, Newark, NJ, United States

Abstract Thermomechanical analysis (TMA) is one of the most important thermal analysis techniques. TMA and thermodilatometry measure some dimension as a function of temperature. The difference between them is that thermodilatometry measures the dimensional changes when negligible load is applied to the sample, while in TMA the load may be significant. TMA for fibers is a frequently used thermal analysis technique, because it can measure the thermal shrinkage and shrinkage force when the temperature is being raised. In this chapter the comparison of shrinkage force is given for drawn, heat set and relaxed fibers, and the origin of shrinkage force is explained. Fibers in most cases are characterized by the coefficient of linear thermal expansion measured in the fiber axis direction. The use of the newest TMA technique, modulated temperature TMA is presented and it is shown how to use the Reversing Dimension Change for characterizing shallow and broad glass transitions.

Thermomechanical analysis (TMA) is an extremely useful thermal-analysis technique with a broad range of practical applications. TMA measures the dimensional changes while the sample is being subjected to a controlled temperature program. It can also measure the forces generated during dimensional changes, in the sample, as a function of temperature. A relatively new TMA technique is modulatedtemperature TMA (MT-TMA). Here, a sinusoidal modulation is overlaid on the linear temperature change. This technique can separate the thermodynamic dimensional quantities (i.e., coefficient of linear thermal expansion, CLTE) from the kinetic effects (usually shrinkage). All major thermo-analytical instrument companies (TA Instruments, PerkinElmer, Hitachi, Mettler-Toledo, Netzsch, Setaram, Instrument Specialists) manufacture TMA instruments. TMA is the simplest thermal analysis (TA) technique to measure the various moduli of plastics. Here, we need to mention a method that grew out of TMA, which is called dynamic mechanical analysis (DMA). DMA is essentially strain or stress-modulated TMA. DMA became an independent TA technique and is usually treated separately from TMA (see Chapter 6: Dynamic mechanical analysis in fiber research). In TMA a probe is put on the sample and the position of this probe is measured by an linear voltage differential transformer (LVDT, see later) as a function of Thermal Analysis of Textiles and Fibers. DOI: https://doi.org/10.1016/B978-0-08-100572-9.00005-7 © 2020 Elsevier Ltd. All rights reserved.

82

Thermal Analysis of Textiles and Fibers

temperature, in most cases, during linear heating. Some weight is put on the probe, and this measurement is called TMA, but there is an almost identical technique called dilatometry. By definition, dilatometry is a zero-force TMA. Unfortunately, there is a mess in the literature regarding the difference between TMA and dilatometry. It is difficult to clear up the reason for this uncertainty. Most often in the published papers, a “mercury dilatometer” is called dilatometer. In this method, a glass capillary is attached to a U-shaped glass tubing as a result of which mercury covers the sample. When the sample is heated, the mercury level increases as the sample expands, and the capillary is used to ensure high-accuracy reading of the mercury level, thus providing a high-accuracy determination of the sample volume. However, the sample is covered by liquid mercury, and knowing the density of mercury, it is difficult to call this a “zero-force” measuring technique. So, the reader is advised to read any paper carefully and understand what the authors mean when they say “dilatometry.” Most often, the temperature dependence of volume determination is called dilatometry, and while TMA is the measurement of the length during expansion. TMA instruments can also perform dilatometry (i.e., volume) measurements: the sample is placed in quartz powder and the expansion is measured. Quartz has very low thermal expansion, but even this expansion is subtracted from the sample curve, thus the volumetric expansion of the sample can be obtained by a TMA instrument. Creep, stress relaxation, and the simple force-measuring determination of the tensile modulus principally belong to the technique of TMA. However, the determination of these physical quantities is usually treated in the discussion of DMA, because it is more advantageous to carry out these measurements with DMA instruments than with TMA instruments. It is true that both TMA and DMA instruments are becoming more and more similar in the sense, that TMA instruments can increasingly perform more DMA measurements and vice versa. This is somewhat similar to digital still cameras performing simple video recordings and video cameras making simple digital still pictures, and if this trend continues, we will have one common instrument that can carry out both types of measurements. So today, TMA instruments can make simple DMA measurements (TAI’s Q400 can make DMA measurements, but the maximum frequency is only 0.5 Hz) or DMA instruments can record TMA curves (see PerkinElmer DMA7 for measuring TMA). The major TMA applications in fiber research are the following: 1. Determination of thermal history of polymeric samples (including fibers). Very often, this is called “fingerprinting.” The sensitivity of TMA instruments is often much higher than the sensitivity of differential scanning calorimetry (DSC) instruments (for this reason, see the description of LVDT later); therefore TMA can be considered a preferred method for carrying out fingerprinting of polymeric samples of any shape (including films and fibers). 2. Determination of end-use properties such as shrinkage, shrinkage force, and coefficient of linear thermal expansion (CLTE). 3. Determination of transition temperatures (melting point Tm, glass transition temperature Tg, etc.). 4. Studying the kinetics of shrinkage and shrinkage force phenomena.

Thermomechanical analysis of fibers

83

Figure 5.1 Schematics of a static force TMA instrument.

The major variables in TMA measurements are the applied load, atmosphere, and time temperature. TMA has a major advantage in fiber research, because the parameters monitored (length change, stress, and time/temperature) are the same as the major variables of fiber processing. This was discussed in detail by Valk (1972) and Addyman and Ogilvie (1979). As mentioned above, TMA is the most sensitive TA technique. Fig. 5.1 shows the schematic diagram of a TMA instrument. The sample is put on the sample platform surrounded by the furnace to heat the sample. The probe is placed on the sample. The probe is usually made of quartz because of its small coefficient of thermal expansion. During the measurement the sample expands and this expansion raises the quartz probe. As mentioned previously, one end of the probe is put on the sample and the position of the other end is measured by a position transformer. The position transformer is a very sensitive device called LVDT (see Fig. 5.2). LVDT stands for “linear voltage differential transformer” or “linear variable differential transformer” (other names are differential transformer, linear variable displacement transformer, and linear variable displacement transducer). Today, it is impossible to figure out which expression was the original name of this device, but it is more important to turn to its structure. The hysteresis of LVDTs is low, and the precision of the distance measurements is very high. LVDT is an electromagnetic transducer, which transforms a position signal into a proportional electrical signal. This electrical signal contains a phase and an amplitude signal (for direction and distance, respectively). An LVDT has three solenoidal coils: the primary coil and two secondaries. The coils are placed end-to-end to each other: the primary coil is in the center and the two secondaries are at the sides. Alternating current flowing through the primary coil generates induction current in the secondary coils. A ferromagnetic core (FC) is attached to the sample probe of

84

Thermal Analysis of Textiles and Fibers

Ferromagnetic core (armature)

Input

Output

VA

Vout = VA – VB

Primary coil

Secondary coil 1

VB

Secondary coil 2

Figure 5.2 The schematics of the LVDT.

the TMA to measure its position. The FC slides along the long axis of the LVDT tube. The voltage of the currents generated in the secondary coils is proportional to the position of the core. The output voltage is the difference between the voltages of the two secondary coils, because the coils are connected. If the FC is exactly in the central position of the LVDT tube (i.e., the same distance from both secondaries), the two generated voltages cancel each other, but when the FC shifts to one side, the voltage in the closer secondary coil increases and the voltage in the other (further) coil decreases. Therefore the total output voltage goes up. In addition, LVDT is direction-sensitive. When the FC moves closer to secondary coil 1, the generated voltage is in phase with the primary voltage. However, when the FC moves toward the secondary coil 2, the generated voltage also increases, but the phase of the generated voltage, this time, is opposite to the primary voltage. The LVDT is a very sensitive device because the side of the FC does not touch any solid material (it is friction free). The maximum sensitivity of an LVDT is around 10 nm. All TMA instruments are purged with some purge gas (in most cases, highpurity dry nitrogen). Almost all significant TA-instrument manufacturers, such as TA Instruments, Perkin-Elmer, Netzsch, Mettler-Toledo, Hitachi, Instrument Specialists, and Setaram, market TMA instruments. The most important variables in TMA measurements are dimension, temperature/time, and applied load. As mentioned above, TMA is one of the best TA techniques for fingerprinting, because the major variables of TMA are the same as the key variables of fiber processing. In fiber TMA, one of the most important difficulties is the preload. A certain minimum load needs to be put on the fiber to straighten it out. If the load is too small, the fiber may not be straightened enough. If the load is too high, it may induce additional orientation to the fiber or break it. If such a problem is encountered, the experiments need to be run at various preload values. For most materials the CLTE is positive during a heating experiment. This is somewhat different for oriented polymers. At least the formal CLTE (that is

Thermomechanical analysis of fibers

85

Figure 5.3 MT-TMA measurement of an as-spun and drawn PET single filament (Draw Ratio, DR 5 4 3 ): the reversing dimension change (diameter of the drawn fiber is 20 nm). The glass transition corresponds to the jump in CLTE. The glass transition temperature (Tg) of the as-spun fiber is lower and sharper than that of the drawn fiber. Solid curve: as-spun fiber, dashed curve: drawn fiber. TA Instruments Q400 TMA (Menczel, 2020).

measured by TMA) can be negative above the glass transition temperature due to shrinkage. MT-TMA can be a big help in these cases. MT-TMA can separate the real (thermodynamic) CLTE of the material from the kinetic dimensional changes such as shrinkage. In the MT-TMA1 experiments performed on TA Instruments’ Q400 TMA, the reversing dimensional change displays the thermodynamic CLTE, while all the kinetic dimensional changes are displayed by the nonreversing dimensional changes (see Fig. 5.3). Therefore if there is any doubt regarding the dimensional change of the fiber, we recommend running an MT-TMA experiment. The aforementioned facts are very important for practical applications. The first MT-TMA-capable TMA instrument was marketed in 1998 by TA Instruments. Before that time, the expression “CLTE of fibers” meant the actual physical quantity measured by the TMA instrument, and this was the sum of the real CLTE of the material (in one direction) and the thermal shrinkage. Today, “CLTE” implies the actual CLTE of the fiber in the z-axis direction. In the Q400 software, this quantity is indicated as “reversing dimensional change.” To make this situation more complicated, we need to mention that CLTE of a fiber indicates a one-dimensional expansion, and we may designate this as CLTEz 1

MTMA (modulated TMA) is the registered trademark of TA Instruments. The name of the technique, by the definition of International Confederation for Thermal Analysis and Calorimetry (ICTAC), is MT-TMA. This is similar to the case of dielectric thermal analysis (DETA)-dielectric analysis (DEA) and dynamic mechanical thermal analysis (DMTA)-DMA. Polymer Labs (presently owned by TA Instruments, i.e., Waters Corporation) introduced DMTA and DETA as trade names, but the ICTACrecognized names (the names of the techniques) are DMA and DEA.

86

Thermal Analysis of Textiles and Fibers

(the CLTE in the fiber long-axis direction). But in various handbooks, CVTE (coefficient of volumetric thermal expansion) is very often given in tables for various polymers (see Brandrup et al., 2003; Wakelin et al., 1960). For isotropic materials CVTE 5 3CLTE. On the contrary, for fibers, this is not true because the CLTEx and CLTEy (the coefficient of linear expansion perpendicular to the long axis of the fiber) have the same value (there is axial symmetry), but the expansion parallel to the fiber axis (CLTEz) has a different value (usually smaller) due to orientation. Also, this is the direction where thermal shrinkage takes place. We need to emphasize that everything you measure in TMA of fibers is CLTEz. The value of the real CLTEz (i.e., the thermodynamic value without shrinkage) is usually less than CLTEx and CLTEy (CLTEx 5 CLTEy 5 CLTE| due to axial symmetry): the system tries to go toward the random coil configuration above the glass transition temperature, and hence the perpendicular components of the CLTE have higher values than the parallel (z-axis) component. Also, CLTE exhibits a much stronger dependence on chain orientation than on crystallinity, as determined by Jaffe (1977). He obtained a reasonable correlation between CLTEz and birefringence for PET yarns. Kimmel (1971) reported that for drawn polyacrylonitrile (PAN) yarns, CLTEz decreases with increasing draw ratio (Fig. 5.4). At the same time, the modulus and tenacity of this fiber increases with increasing CLTEz, as expected. Porter et al. (1975) obtained negative values (i.e., shrinkage) for CLTEz in ultraoriented polyethylene fibers. These authors also indicated that CLTEz of fibers should be controlled for composite applications: in such cases, similar expansion properties of the fiber and the matrix will reduce thermal stresses at the fiber matrix interfaces, thus reducing the probability of failure in the composites. For end-use potential of a given fiber, thermal shrinkage is one of the most important properties (setting safe ironing temperature of textiles, selecting a fiber for reinforcing tire cords, etc.). Few fiber properties reflect the process history or

Figure 5.4 The axial CLTE of a typical acrylic yarn as a function of draw ratio. Source: From Kimmel, R.M., 1971. Fiber Soc. Lect. Reprinted with permission of Kimmel, R.M.

Thermomechanical analysis of fibers

87

the structural state of the fiber as completely as shrinkage (shrinkage is a very sensitive property for structural fingerprinting). When fiber dimension of a drawn semicrystalline fiber plotted as a function of temperature at zero load four different regions can be distinguished in this diagram: (1) The region up to Tg is characterized by reversible thermal expansion, but small amounts of shrinkage can occur in this region if the fiber contains some residual solvent or moisture, and these leave the fiber, especially at high draw ratios. In this region, dSh/dT is small and constant. The thermodynamic CLTE is close to the value measured by actual TMA measurement. (2) The region of the glass transition. Rapid irreversible shrinkage takes place in this region due to relaxation of relatively free oriented amorphous chain segments. Of course, the peak in the dSh/dT curve corresponds to the maximum shrinkage rate. (3) In the temperature region between Tg and start of the melting process or start of reorganization, additional shrinkage takes place, because the mobility of some former crystalline segments suddenly increases due to recrystallization or melting. Also, the so-called entropy shrinkage that was started in Region 2 is continuing. This type of shrinkage depends mostly on process history, and the peaks in the dSh/dT curve in this region do correspond to heat treatment (DSC “premelting”) temperatures (annealing peaks) (DSC is differential scanning calorimetry). (4) Prior to fiber failure, rapid shrinkage takes place in this region due to melting of the crystalline regions. The described fiber shrinkage mechanism is based on the two-phase model of semicrystalline polymers. It needs to be mentioned that not all the described shrinkage regions will be seen for every fiber: the different fiber-process histories will influence the shrinkage behavior of the specific fiber. The views on the shrinkage behavior of fibers are vastly different in the literature (Bhatt and Bell, 1976; Bosley, 1967; Pokrivskaya and Uterskii, 1972; Prevorsek et al., 1974; Ribnick, 1969a,b; Samuels, 1974; Shishoo and Bergh, 1977; Wilson, 1974). Shrinkage processes are irreversible as proven by Berndt and Heidemann (1977). These authors studied the behavior of nylon 6 and poly(ethylene terephthalate) (PET). They heated the fibers to temperatures Th beyond Tg, and thus eliminated part of the shrinkage. They then cooled the samples and reheated them, and they observed that shrinkage recommences beyond Th. This means that only the length changes associated with the thermodynamic CLTEs are reversible, whereas shrinkage itself is irreversible. Shrinkage behavior of acrylic fibers was studied in detail by Kimmel (1971). He found good correlation between the shrinkage curves and processing conditions as shown in Fig. 5.5. The mentioned author also studied correlation between the melting data of as-received acrylic yarns with varying comonomer content as determined by DSC and TMA: the peak temperature of the DSC melting peaks was plotted as a function of position of the high temperature dST/dT peak from TMA shrinkage experiments. A good correlation was obtained (Fig. 5.6). As described later in this chapter, TMA is a good complimentary technique to DSC, because volume is a thermodynamic analogue of the enthalpy, and very often, the change in volume is more pronounced than the change in enthalphy.

88

Thermal Analysis of Textiles and Fibers

Shrinkage (%)

(A)

30

20

Drawn Heat-set

10

Relaxed

0 0

50

100 150 200 Temperature (°C)

250

300

(B)

Shrinkage rate (%/ºC)

0.4 Relaxed

0.3

Heat-set

0.2 Drawn

0.1

0 0

50

100

150

200

250

300

Temperature (°C)

Figure 5.5 (A) TMA curves of drawn, heat-set and relaxed acrylic yarn. (B) Shrinkage rate curves (dSh/dT) of the acrylic yarns shown in (A). Source: From Kimmel, R.M., 1971. Fiber Soc. Lect. Reprinted with permission of Kimmel, R.M.

Additional shrinkage takes place for fibers spun at very spinning high speeds. This shows the presence of oriented structures stabilized by crystalline units, but the increase of crystallinity, in general, tends to decrease shrinkage and oppositely, increased orientation tends to increase shrinkage (Samuels, 1974). Fig. 5.7 shows shrinkage curves of as-spun PET yarns for various wind-up speeds (Heuvel and Huisman, 1978). When the spinning speed (that is proportional to melt stress) increases, the shrinkage, just above Tg, first increases and then decreases. It is easy to explain the initial shrinkage increase, which is due increase of orientation of noncrystalline regions. It is more difficult to explain the later decrease, but it can be suggested that it is due to the onset of crystallization during spinning. When the fiber is heat-treated during processing (this process is called “annealing”) at temperatures below the major melting endotherm, another smaller

Figure 5.6 Correlation of the melting point of acrylic copolymer yarns as determined by TMA and DSC. The TMA temperatures are the temperatures at maximum shrinkage rates of 10 C/min, the DSC temperatures are the peak temperatures of melting at 5 C/min heating rate. DuPont (TA Instruments) 990-941 TMA and 990-910 DSC. The numbers in parentheses indicate the acrylonitrile content (%) in the copolymers. Source: From Kimmel, R.M., 1971. Fiber Soc. Lect. Reprinted with Permission of Kimmel, R.M.

Figure 5.7 TMA shrinkage curves of as-spun PET yarns for different wind-up speeds. PerkinElmer TMS-1 TMA instrument, 0.04 GN/tex applied load, 20 C/min heating rate. Source: From Heuvel, M., Huisman, R., 1978. J. Appl. Polym. Sci. 22, 2219, Fig. 11. Reprinted with permission of Wiley.

90

Thermal Analysis of Textiles and Fibers

magnitude endothermic peak is displayed in the DSC curve. The magnitude of this peak increases with time of annealing. By some reason, Berndt and Bossmann (1976) called the peak temperature of this small endotherm Teff. Teff is 5 C 10 C lower than the heat-treatment temperature. Heidemann and Berndt (1976) also defined an effective stress σeff for TMA experiments: σeff is the stress under which the fiber will be at its original length despite its shrinkage, that is, this is the stress necessary to balance the shrinkage force generated at Teff. It is not difficult to measure σeff: the preloading on the fiber must be increased until an SF curve with zero slope at Teff is achieved. The value of shrinkage force (SF) in the zero-slope region will determine σeff. Above Teff, the SF will suddenly decrease, because the structure retaining the frozen-in stress was destroyed. Therefore with such an experiment, both Teff and σeff can be simultaneously determined. Later, Berndt and Heinemann (1980) introduced a modified method for simultaneous determination of σeff and Teff. In this method, small strain adjustments are made to the fiber during heating in the SF instrument (this instrument was Textechno Thermofil) so that the stress relaxation and the shrinkage force development (i.e., stress retardation) are balanced. This method balances the relaxation and retardation processes: the SF will have a maximum value at Teff determining the value of σeff. These parameters may be used to predict the behavior of fibers when they are produced differently, and the intention is to obtain fibers of similar properties or determine causes of fiber nonuniformity. Kimmel (1971) investigated the effect of heat setting and relaxation on the SF curve of drawn acrylic yarns. The results are shown in Fig. 5.8. Fig. 5.9 shows the

Figure 5.8 Shrinkage force curves of drawn, heat-set, and relaxed acrylic yarns. These measurements were carried out on a homemade shrinkage force analyzer by Kimmel (1971). The instrument utilized an Instron tensile tester and a DuPont (presently TA Instruments) 900 temperature programmer controlling the heating of a clamshell-type oven with a 0.5-in. bore and a 6-in. gauge length. The heating rate was 10 C/min. Source: From Kimmel, R.M., 1971. Fiber Soc. Lect. Reprinted with permission of Kimmel, R.M.

Thermomechanical analysis of fibers

91

Figure 5.9 Shrinkage force curves of heat setting temperature on the shrinkage force curves of drawn acrylic yarns. Source: From Kimmel, R.M., 1971. Fiber Soc. Lect. Reprinted with permission of Kimmel, R.M.

Figure 5.10 Shrinkage force curves for PET yarns as a function of temperature for various pretreatment length changes. Source: From Heidemann, G., Berndt, H.J., 1974. Chemiefasern Text.-Amwendungstech./ Text. Ind. 24 (1), 46, Fig. 9. Reprinted with permission of Deutscher Fachverlag GmbH.

sensitivity effect of processing temperature of the shrinkage force response of a drawn acrylic fiber. Heidemann and Berndt (1974) studied how the processing strain and stress influence the shrinkage force for drawn acrylic yarns. They also studied the kinetics of

92

Thermal Analysis of Textiles and Fibers

Figure 5.11 Shrinkage force curves of PET yarns as a function of temperature for two different pretreatment tensions and temperatures. The heating rate was 12 C/min. Source: From Heidemann, G., Berndt, H.J., 1974. Chemiefasern Text.-Amwendungstech./ Text. Ind. 24 (1), 46, Fig. 10. Reprinted with permission of Deutscher Fachverlag GmbH.

Figure 5.12 The maximum shrinkage force as a function of draw ratio for [poly(ethylene terephthalate)] PET yarns. The shrinkage force was measured with a “Statigraph” strength tester connected to a temperature control device. 0.5 p/tex starting stress, 200 mm gauge length. Source: From Hoffrichter, S., 1973. Faserforsch. Textiltech. 24 (7), 289, Fig. 14. Reprinted with permission of AKS Demie-Verlag Berlin.

shrinkage force buildup and decay. Similar to shrinkage, shrinkage force tends to increase with increasing orientation and decreasing crystallinity. The relationship between shrinkage and shrinkage force is complex, it is not easy to find some correlation between them (see Figs. 5.10 and 5.11). Fig. 5.12 shows the expected increase of shrinkage force with increasing draw ratio (i.e., increasing orientation) for a series of polyester yarns (Hoffrichter, 1973).

Thermomechanical analysis of fibers

93

Figure 5.13 The various components of the dimension changes in an MT-TMA experiment of a drawn PET single filament using a TA Instruments Q400 TMA instrument with temperature modulation. (Menczel, 2020).

Lately, Menczel carried out a series of MT-TMA measurements on drawn semicrystalline fibers. An example of these results is shown in Fig. 5.13. From the curves, it is clear that shrinkage is the major component of the dimensional change, the total dimension change curve is almost identical to the nonreversing dimension change. At the same time, the reversing component of the dimension change is very useful, but much smaller in magnitude (watch the scales!), since it shows the glass transition of the drawn fiber much clearer than the DSC measurements (compare to DSC curves in Chapter 3: Differential scanning calorimetry in fiber research). MT-TMA measurements of Menczel on Dyneema ultrahigh modulus polyethylene fibers gave a Tg of B 2 20 C in good agreement with the heat capacity measurements of Gaur and Wunderlich (Gaur and Wunderlich, 1980).

References Addyman, L., Ogilvie, G.D., 1979. Br. Polym. J. 11 (3), 15. Berndt, H.J., Bossmann, A., 1976. Polymer 17, 241. Berndt, H.J., Heidemann, G., 1977. Melliand Textilber. 1, 83. Berndt, H.J., Heinemann, G., 1980. Proc. Int. Conf. Therm. Anal. 6th 1, 345. Bhatt, G.M., Bell, J.P., 1976. J. Polym. Sci., Phys. 14, 575. Bosley, D.E., 1967. J. Polym. Sci., C 20, 77.

94

Thermal Analysis of Textiles and Fibers

Brandrup, J., Immergut, E.H., Grulke, E.A. (Eds.), 2003. Polymer Handbook. fourth ed. Wiley. Gaur, U., Wunderlich, B., 1980. Macromol. 13, 445. Heidemann, G., Berndt, H.J., 1974. Chemiefasern Text.-Amwendungstech./Text. Ind. 24 (1), 46. Heidemann, G., Berndt, H.J., 1976. Melliand Textilber. (Engl. Ed.) 6, 485. Heuvel, M., Huisman, R., 1978. J. Appl. Polym. Sci. 22, 2219. Hoffrichter, S., 1973. Faserforsch. Textiltech. 24 (7), 289. Jaffe, M., 1977. Therm. Anal. Symp., Am. Phys. Soc. Prepr. 1977, 1. Jaffe, M., Menczel, J.D., Bessey, W., 1997. Chapter 7: Fibers. In: Turi, E.A. (Ed.), Thermal Characterization of Polymeric Materials. Academic Press, San Diego, CA, pp. 1767 1954. Kimmel, R.M., 1971. Fiber Soc. Lect. Menczel, J.D., 2020. To be published. Pokrivskaya, L.V., Uterskii, L.E., 1972. Khim. Volokna 14 (2), 10. Porter, R.S., Weeks, N.E., Capiati, N.J., Krzewski, R.J., 1975. J. Therm. Anal. 8, 547. Prevorsek, D.C., Tirpak, G.A., Harget, P.J., Reimschuessel, A.C., 1974. Macromol. Sci., Phys B9 (4), 733. Ribnick, A., 1969a. Text. Res. J. 39, 428. Ribnick, A., 1969b. Text. Res. J. 39, 742. Samuels, R.J., 1974. Structured Polymer Properties. Wiley Interscience, New York. Shishoo, R., Bergh, K.M., 1977. Text. Res. J. 47, 56. Valk, G., 1972. Lenzinger Ber. 33, 1. Wakelin, J.H., Sutherland, A., Beck, L.R., 1960. J. Polym. Sci. 42, 278. Wilson, M.P.W., 1974. Polymer 15, 277.

Further reading Valk, G., Berndt, H.J., Heidemann, G., 1971. Chemiefasern 5, 1.

Dynamic mechanical analysis (DMA) in fiber research

6

Joseph D. Menczel Thermal Measurements LLC, Fort Worth, TX, United States

Abstract Dynamic mechanical analysis (DMA) is essentially a strain or stress-modulated thermomechanical analysis (TMA). It is widely used for determining mechanical properties of polymers and fibers and measuring the glass transition temperature and other second-order transitions (α, β, etc.) in polymers and fibers. The technique of dynamic mechanical analysis and its peculiarities for fiber measurements are described. It is explained how DMA can help to determine the necessary drawing conditions for preparing ultraoriented fibers [poly(vinyl alcohol)].

Before presenting peculiarities of dynamic mechanical analysis (DMA) measurements of fibers and films, let us briefly describe some basic aspects of preparation of synthetic fibers. Fibers from synthetic polymers are made by spinning, which means that the molten (or dissolved) polymer is continuously pushed through narrow holes. The filaments formed this way are cooled rapidly, and this process leads to the formation of low crystallinity fibers, called as-spun fibers. Since these fibers usually are not oriented (they are with little or no crystallinity), they are stretched significantly at some optimum temperature to develop the necessary orientation and crystallinity so that their modulus and tenacity increase. These fibers are then called drawn fibers and are characterized by the draw ratio, which is the ratio of the final crosssectional area over the initial cross-sectional area. Special areas of fiber preparation are gel-spinning and gel-drawing (Pennings, 1967; Smook and Pennings, 1984). When the polymer chains are disentangled from each other in a dilute solution, this dilute solution can be spun into a coagulation bath, and this way a solid gel is formed. Then, these gel-spun fibers can be drawn to very high draw ratios (since the individual polymer chains are disentangled from each other) to impart extremely high orientation. This way, ultrahigh molecular mass polyethylene (UHMMPE) can be drawn up to a draw ratio of 300 3 , and polypropylene to 57 3 . Extremely high tenacities can be achieved with gel-spun and gel-drawn fibers. A certain number of single filaments (with different name, monofilaments) need to be put together and then twisted in order to increase the friction between the monofilaments. The name of these interlocked fibers is “yarn.” The mechanical properties of yarns are not simply the sum of the mechanical properties of the Thermal Analysis of Textiles and Fibers. DOI: https://doi.org/10.1016/B978-0-08-100572-9.00006-9 © 2020 Elsevier Ltd. All rights reserved.

96

Thermal Analysis of Textiles and Fibers

constituting single fibers. Thus, when reporting the results, it must always be mentioned whether the measurements were made on single fibers or yarns. A special method, called zone drawing
zone annealing, was developed by Kunugi (see Kunugi et al., 1986, 1988, 1991) to prepare high modulus and hightenacity fibers and films from various polymers [polyethylene, polypropylene, poly (ethylene terephthalate) (PET), poly(ether ether ketone) (PEEK), nylons, poly(vinyl alcohol) (PVA)]. From sample manipulation point of view, it is not easy to perform DSC measurements on fibers. Experienced eyes are needed to recognize the glass transition of highly drawn fibers, since the heat capacity increase of drawn fibers at the glass transition is broad and shallow. This is due to the presence of the rigid amorphous fraction and the oriented amorphous fraction (see Menczel and Wunderlich, 1981). Similarly, TMA does not provide an easily noticeable break in the V 5 f(T) of ‘ 5 f (T) curves to clearly characterize the glass transition (V is the volume or specific volume and ‘ is the sample length). Therefore DMA is very popular and it is the most frequently used technique in the fiber and film technology for characterization of the glass transition, because the glass transition appears as a peak in the tan δ versus temperature or the loss modulus versus temperature curve. In addition, one measurement will provide the temperature dependence of the modulus, the glass transition temperature, and characterize possible second-order transitions of the material. Details of DMA measurements of fibers and films can be found in several references (see e.g., Murayama, 1978; Jaffe et al., 1997; Menczel et al., 1997b). Here, we briefly summarize the most important aspects of this special area of DMA measurements. DMA measurements of fibers and thin films require special attention due to the unusual mechanical properties of fibers and films originating from their unique geometry. The special geometry of fibers and films allows DMA measurements in tension only: due to the small diameter of polymeric fibers and small thickness of polymeric thin films (and also often elastomeric thick films), these samples are unable to support significant levels of compressive stress. Therefore some static stress has to be applied to the fiber or the film sample that is greater than the peak level of the stress attained in the dynamic oscillation. If such static stress is absent during the measurements, sample buckling will occur, as shown in Fig. 6.1, and the lower part of the stress signal will be truncated (Grehlinger and Kraft, 1988a,b). The just-mentioned additional static force is called “pretension.” As can be seen in Fig. 6.1, the pretension raises the height level of the sine stress wave, so the originally zero level of the midpoint of the sine wave will get a positive value. The lowest level of the pretension must be higher than the half height of the peak-to-peak value of the dynamic force to avoid sample buckling. However, it is not always easy to determine the highest possible level of the pretension, because too high a pretension can lead to permanent changes in the sample structure, for example, it can increase the orientation of the fiber. Grehlinger and Kraft (1988a,b) described the effect of pretension on the storage modulus of fibers. Another problem with highly oriented (i.e., drawn) fibers or films is the onset of shrinkage as the temperature is being raised during the DMA measurement.

Dynamic mechanical analysis (DMA) in fiber research

97

Figure 6.1 A figure explaining the necessity of using a pretension in DMA measurements of fibers and thin films: sample buckling will take place below a certain pretension level. Source: From Grehlinger, M., Kraft, M., 1988a. J. Plast. Film Sheeting 4, 318. Reprinted with permission of Technomic Publishing Co., Inc.

Creep may also take place. The presence of several physical processes taking place during the heating of the fiber (change of the modulus with temperature, shrinkage, creep) requires monitoring and continuous adjustment of the pretension level. Therefore most commercial instruments do have an automated mechanism to deal with this problem, and this mechanism is called “autotension.” The pretension can be selected as follows: 1. It can stay constant during the whole run. 2. It can be predicted from the change of the dynamic force with temperature. 3. It can be predetermined using the temperature dependence of the storage modulus of the fiber.

Static stress
strain measurements are usually needed for estimating the pretension level. In these measurements, one pulls the fiber or the film in tension mode at a constant rate and then plots the resulting stress or force as a function of strain. The pretension needs to be selected such that the total stress during the DMA measurements falls onto the linear starting portion of the stress
strain curve (Fig. 6.2). Obviously, the autotension mode will influence the DMA curve, since it may lead to different levels of shrinkage and creep. Almost all presently available commercial DMA instruments have special fiber and film probes. Instruments specifically designed to run fibers are commercially available (e.g., Rheometrics RSA).

98

Thermal Analysis of Textiles and Fibers

Figure 6.2 The stress
strain curve in tensile mode.

In DMA measurements of fibers and films, one can get information about the various transitions using the temperature dependence of the storage and loss moduli. In order of decreasing temperatures, these transitions are as follows: G

G

G

G

G

The so-called αc-transition is the highest temperature transition. This is a relaxation process taking place in the crystallites, therefore at temperatures below the melting point, but higher than the glass transition temperature. Cold crystallization (see Section 3.1, in Chapter 3: Differential scanning calorimetry in fiber research), when crystallization takes place during the heating of a quenched crystallizable amorphous polymer, for example, PET and poly(ethylene-2,6-naphthalene dicarboxylate) (PEN). In such cases when the temperature is being scanned (increased), the slow modulus decrease is followed by a sudden drop (of 2
3 orders of magnitude) due to the glass transition. Then at somewhat higher temperatures, a sudden modulus increase takes place due to cold crystallization, since the modulus of a semicrystalline polymer is always higher than the modulus of a similar amorphous polymer. (see Fig. 6.4). The amorphous α-relaxation (or αa-relaxation) which is essentially the glass transition. The various second-order transitions taking place in either the crystalline or the amorphous phase (or in both). DMA measurements can provide information about the orientation of the fiber or film through the magnitude of the storage modulus or by the glass transition temperature and width of the glass transition.

An interesting area is the effect of crystallinity and orientation of the fiber on the glass transition temperature. Several research groups noticed that Tg of the fibers at first increases as the crystallinity increases, then with further crystallinity increase, Tg starts decreasing (see e.g., Dumbleton and Murayama, 1967; Dumbleton et al., 1968; Illers and Breuer, 1963; Thompson and Woods, 1956). This phenomenon was explained by the change in the number and size of the crystallites in the fiber. The effect of orientation is more straightforward: in most cases, Tg of the fiber increases with increasing orientation. The reason for this increase is the decrease of the free volume with increasing orientation. However, the mentioned dependence can be complex, if the crystallinity of the fibers changes simultaneously

Dynamic mechanical analysis (DMA) in fiber research

99

Figure 6.3 (A) Storage modulus of a PET yarn from fibers drawn to various draw ratios. The modulus increases with increasing draw ratio. Rheovibron viscoelastometer, test frequency 11 Hz, heating rate 1 C/min. (B) Tan δ (5E00 /E0 ) of PET yarn prepared from fibers drawn to various ratios. As the orientation (i.e., draw ratio) and crystallinity increases, the glass transition is becoming broader and shallower. Rheovibron viscoelastometer, test frequency 11 Hz, heating rate 1 C/min. Source: (A and B) From Miller, R.W., Murayama, T., 1984. J. Appl. Polym. Sci. 29, 933 (Miller and Murayama, 1984). Reprinted with permission of Wiley.

with orientation. Also, the width of the glass transition changes significantly with orientation. Fig. 6.3A shows how the storage and loss moduli of a PET yarn change as the draw ratio of the fibers constructing the yarn increases. It is obvious that the

100

Thermal Analysis of Textiles and Fibers

1011 DR = 5.4 1010

E′ (—) (Pa)

109

DR = 2.3 As-spun fiber

108 107 106 105 0.0

50.0

100.0 150.0 Temperature (ºC)

200.0

250.0

Figure 6.4 Tensile storage modulus of as-spun and drawn PEN fibers as a function of temperature. Rheometrics RSA2 DMA, heating rate 1 C/min. Source: From Saw, C.K., Menczel, J., Choe, E.W., Highes, O.R., 1997. SPE ANTEC ’97, April 27
May 2, 1997, Toronto, vol. 2, Materials, p. 916. (Saw et al., 1997). Reprinted with permission of Society of Plastics Engineers.

storage modulus increases with increasing draw ratio (since the orientation of these fibers increases with increasing draw ratio). The orientation and the crystallinity also have an effect on the glass transition as can be seen on the loss modulus versus temperature curves in Fig. 6.3B: the yarn of the as-spun fibers has a sharp glass transition, because it is not oriented and close to being amorphous. Fig. 6.4 shows the dynamic mechanical curves of as-spun and drawn PEN monofilaments. The as-spun fiber, which is unoriented and amorphous, exhibits a slow modulus decrease with temperature, then, beyond 100 C, the modulus suddenly drops due to the glass transition of the polymer. This modulus decrease is of 2
3 orders of magnitude. Then, the modulus increases again due to cold crystallization: a crystalline polymer always has higher modulus than an amorphous one (of the same polymer). Finally, with further temperature increase, the storage modulus decreases again, this time due to melting of the PEN crystallites. Of course, there is no modulus increase in the drawn fibers due to cold crystallization, because these fibers are crystalline, so no cold crystallization can take place. The modulus increase of the as-spun fiber at the glass transition due to cold crystallization is ca.1 order of magnitude. This picture is typical of semicrystalline polymers that can be quenched to an amorphous state. Garrett and Grubb (1988) described an excellent example of how a crystalline relaxation can affect the drawing conditions of a fiber. Working with highly drawn gel-spun PVA fibers, they observed an intense crystalline αc-relaxation. The temperature of this relaxation (which was determined as the peak temperature of the tan δ vs T curves) depended on the draw ratio of the fiber. At a draw ratio of 1

Dynamic mechanical analysis (DMA) in fiber research

(A)

101

(B)

12.0

λ = 2x 11.5

λ = 38x λ = 15x

λ = 5x

λ = 8.5x 0.2

10.5 Tan δ

Log(E′) (dyne/cm2)

11.0

λ = 2x 10.0

0.0

λ = 15x

9.5

9.0

8.5 –150 –100 –50

λ = 8.5x

λ = 38x

0

50

100 150 200

Temperature (ºC)

250

–150 –100 –50

0

50

100 150 200 250

Temperature (ºC)

Figure 6.5 (A) Storage modulus of PVA fibers gel drawn to various ratios. The draw ratio is indicated at each curve. Rheovibron DDV-II-C viscoelastometer, test frequency 3.5 Hz, heating rate 3 C/min. (B) Tan delta of PVA fibers gel drawn to various ratios. The draw ratio is indicated at each curve. The temperature of the αc-relaxation (the higher temperature relaxation on each curve) shifts to higher temperatures with increasing draw ratio moving beyond the melting point of PVA at DR . 38. Rheovibron DDV-II-C viscoelastometer, test frequency 3.5 Hz, heating rate 3 C/min. Source: (A and B) From Garrett, P.D., Grubb, D.T., 1988. J. Polym. Sci., B: Polym. Phys. 26, 2509, Figs. 4 and 5. Reprinted with permission of Wiley.

(i.e., for undrawn fiber) Tα 5 160 C, and it moves up to 220 C at a draw ratio of 32 3 (see Fig. 6.5). It was suggested that at very high draw ratios, the temperature of the αc-relaxation moves beyond the melting point of PVA making the gel drawing impossible. Crystal-to-crystal transitions can sometimes appear on the DMA recordings. Fig. 6.6 shows the storage modulus versus temperature curves for an as-spun and drawn fiber prepared from poly(2-methylpentamethylene terephthalamide) (Nylon M5T, compare with the DSC curves in Section 3.1, in Chapter 3: Differential scanning calorimetry (DSC) in fiber research) (see Menczel et al., 1996; Menczel, 2020). The storage modulus of the drawn fiber, as expected, exhibits slow decrease with temperature, then a sudden drop of B1 order of magnitude occurs due to the glass transition at ca.150 C. Then further slow modulus decrease can be seen due to the slow melting process. The behavior of the as-spun fiber is drastically different. At the glass transition (since it is amorphous, Tg of the as-spun fiber is somewhat lower than Tg of the semicrystalline drawn fiber), the storage modulus drops B3 orders of magnitude, then it suddenly increases due to cold crystallization. Cold crystallization in this case creates a semicrystalline fiber consisting

102

Thermal Analysis of Textiles and Fibers

1012 Drawn fiber (DR=5.2×) 1011

E′ (—) (Pa)

1010

As-spun fiber

109

108

107

106 –150 –100

–50

0

50

100

150

200

250

300

Temperature (°C)

Figure 6.6 Tensile storage modulus of as-spun and drawn Nylon M5T fibers as a function of temperature. Rheometrics RSA2 DMA. Source: From Menczel, J.D., Jaffe, M., Saw, C.K., Bruno, T.P., 1996. J. Therm. Anal. 46, 753. Menczel, J.D., 2019, to be published. Reprinted with permission of Springer.

of amorphous and crystalline regions (we denote this crystal form as A). Then the modulus decreases again due to melting of crystal form A, and it increases due to recrystallization into crystal form B. Then crystal form B melts, and crystal form C starts developing. This is shown on the DMA curve as a modulus decrease and subsequent increase. Finally, crystal form C melts into a smectic phase, and the smectic phase melts into isotropic melt as indicated by the final drop in the storage modulus. The magnitude of the drop in the tensile modulus helped to conclude that the crystal-to-crystal transitions proceed through melting, that is, the transition details look like crystal A!melt!crystal B and crystal B!melt!crystal C. DMA measurements on thin films are similar, but somewhat more complex than those performed on fibers: there is an additional variable in these measurements (the angle between the chain orientation of the film and the strain direction). Changing this angle can drastically change the results. Fig. 6.7 (Menczel et al., 1997a) shows the results for the Vectran liquid crystalline polymer film: it is clearly seen that the modulus of the film is the highest when the strain direction is parallel to the chain orientation. Common sense tells us that the storage tensile modulus of polymeric films must be higher in the draw direction than in the direction perpendicular to the draw (i.e., E|| . E\). This is often the case, but not always. When these moduli (i.e., E|| and E\) are compared for cold drawn polyethylene, polypropylene, polyoxymethylene, poly(ethylene oxide), and Nylon 66 films, E|| is indeed higher than E\. However, when these samples are annealed, the resulting E|| versus T and .E\ versus T

Dynamic mechanical analysis (DMA) in fiber research

103

1011

Vectran film DMA 0 ºC

E′ (Δ) (Pa)

1010

30ºC

109

108

60ºC 90ºC

107 0

50

100

150

200

250

Temperature (°C)

Figure 6.7 Dynamic tensile modulus of Vectran liquid crystalline film samples as a function of temperature. The number at each curve indicates the angle between the strain direction and the direction of chain orientation. Rheometrics RSA2 DMA. Source: From Menczel, J.D., Collins, G.L., Saw, C.K., 1997a. J. Therm. Anal. 49, 201. Reprinted with permission of Springer.

curves will intercept above the glass transition temperature, beyond the glass transition temperature E|| , E\ (Takayanagi et al., 1966; Gupta and Ward, 1967, 1968; Dumbleton and Murayama, 1970; for detailed description see Menczel et al., 1997b).

References Dumbleton, J.H., Murayama, T., 1967. Kolloid Z. Z. Polym. 220, 41. Dumbleton, J.H., Murayama, T., 1970. Kolloid Z. Z. Polym. 238 (1
2), 410. Dumbleton, J.H., Murayama, T., Bell, J.P., 1968. Kolloid Z. Z. Polym. 228, 54. Garrett, P.D., Grubb, D.T., 1988. J. Polym. Sci., B: Polym. Phys. 26, 2509. Grehlinger, M., Kraft, M., 1988a. J. Plast. Film Sheeting 4, 318. Grehlinger, M., Kraft, M., 1988b. SPE ANTEC 88, Atlanta, GA, 1988, p. 1185. Gupta, V.B., Ward, I.M., 1967. J. Macromol. Sci., Phys. B1 (2), 373. Gupta, V.B., Ward, I.M., 1968. J. Macromol. Sci., Phys. B2 (1), 89. Illers, K.H., Breuer, H., 1963. J. Colloid Sci. 18, 1. Jaffe, M., Menczel, J.D., Bessey, W.E., 1997. Chapter 7, “Fibers”, Thermal Characterization of Polymeric Materials, second ed. Academic Press, pp. 1767
1954, 1997. Kunugi, T., Ichenose, T., Suzuki, A., 1986. J. Appl. Polym. Sci. 31, 429. Kunugi, T., Oomori, S., Mikami, S., 1988. Polymer 29, 814. Kunugi, T., Hayakawa, T., Mizushima, A., 1991. Polymer 32, 808. Menczel, J.D., 2020, to be published.

104

Thermal Analysis of Textiles and Fibers

Menczel, J.D., Wunderlich, B., 1981. J. Polym. Sci., Polym. Lett. Ed. 19, 261 (1981). Menczel, J.D., Jaffe, M., Saw, C.K., Bruno, T.P., 1996. J. Therm. Anal. 46, 753. Menczel, J.D., Collins, G.L., Saw, C.K., 1997a. J. Therm. Anal. 49, 201. Menczel, J.D., Jaffe, M., Bessey, W.E., 1997b. Chapter 8, “Films”, Thermal Characterization of Polymeric Materials, second ed. Academic Press, pp. 1955
2089, 1997. Miller, R.W., Murayama, T., 1984. J. Appl. Polym. Sci. 29, 933. Murayama, T., 1978. Dynamic Mechanical Analysis of Polymeric Material. Elsevier, Amsterdam. Pennings, A.J., 1967. Cryst. Growth, Proc. Int. Conf. Boston, 1966, p. 389. Saw, C.K., Menczel, J., Choe, E.W., Highes, O.R., 1997. SPE ANTEC ’97, April 27
May 2, 1997, Toronto, vol. 2, Materials, p. 916. Smook, J., Pennings, J., 1984. Colloid Polym. Sci. 262, 712. Takayanagi, M., Imada, K., Kajiyama, T., 1966. J. Polym. Sci., C 15, 263. Thompson, A.B., Woods, D.W., 1956. Trans. Faraday Soc. 52, 1383.

Thermal analysis of natural fibers

7

Ye Xue1,2, Wenbing Hu3 and Xiao Hu1,2,4 1 Department of Physics and Astronomy, Rowan University, Glassboro, NJ, United States, 2 Department of Biomedical Engineering, Rowan University, Glassboro, NJ, United States, 3 Department of Polymer Science and Engineering, Nanjing University, Nanjing, P.R. China, 4Department of Molecular & Cellular Biosciences, Rowan University, Glassboro, NJ, United States

Abstract Thermal analysis is a critical, analytical, and characterization tool in the field of materials sciences. Specific thermal properties of synthetic polymers and biomaterials with different phases and morphology can be determined through this technique. Traditional thermal analysis techniques include differential scanning calorimetry (DSC), thermogravimetric analysis, thermomechanical analysis, and dynamic mechanic analysis. Among these techniques, DSC is the core thermal analysis technique. It was further developed into modulated-temperature DSC, quasi-isothermal DSC as well as fast DSC. The various DSC techniques and methods characterize the temperatures and heats and/or specific heat capacity changes at the thermodynamic and kinetic transitions of different materials such as low-molecular-mass substances, amorphous and semicrystalline synthetic polymers, and also biopolymers. Moreover, DSC can also help monitor the structural changes of polymers during the heating, cooling, and isothermal measurements. In addition, the calculation of the reversing and nonreversing heat flow can help separate the various transitions.

7.1

Introduction

Thermal analysis is a critical, analytical, and characterization tool in the field of materials sciences (Hu et al., 2006; Huang and Yang, 2005; Menczel and Prime, 2014; Wang et al., 2010). Specific thermal properties of synthetic polymers and biomaterials with different phases and morphology can be determined through this technique. Traditional thermal analysis techniques include differential scanning calorimetry (DSC), thermogravimetric analysis (TGA), thermomechanical analysis (TMA), and dynamic mechanic analysis (DMA) (Wang et al., 2003, 2010; Wilkie, 1999; Xie et al., 2008). Among these techniques, DSC is the core thermal analysis technique. It was further developed into modulated-temperature DSC (MT-DSC), quasi-isothermal DSC as well as fast DSC (F-DSC) (Iervolino et al., 2011; Mathot et al., 2011; Schawe, 2014; Van Herwaarden et al., 2011). The various DSC Thermal Analysis of Textiles and Fibers. DOI: https://doi.org/10.1016/B978-0-08-100572-9.00007-0 © 2020 Elsevier Ltd. All rights reserved.

106

Thermal Analysis of Textiles and Fibers

techniques and methods characterize the temperatures and heats and/or specific heat capacity changes at the thermodynamic and kinetic transitions of different materials such as low-molecular-mass substances, amorphous and semicrystalline synthetic polymers, and also biopolymers. Moreover, DSC can also help monitor the structural changes of polymers during the heating, cooling, and isothermal measurements. In addition, the calculation of the reversing and nonreversing heat flow can help separate the various transitions (Hu et al., 2007, 2008). Natural fibers are environment friendly and can often be recycled. That is why they are widely used in our daily life at the present stage of technological development. They are raw material sources for clothes, textiles, sutures, or novel biomaterials due to their acceptable tensile strength, extensibility, toughness, chemical stability, electrical insulation properties, and biocompatibility (Ma et al., 2005; Scheller et al., 2001; Tajvidi et al., 2006; Thygesen et al., 2005). However, most natural fibers are composed of proteins and polysaccharides and their structures are much more complicated than those of synthetic polymers. Therefore it is essential to characterize natural fibers, such as silk, resilin, elastins, hemicellulose, lignin, collagens, and cellulose, through novel thermal analysis techniques and methods. This will provide a better understanding of their hierarchical structures and morphologies as well as their physical properties and self-assembly mechanisms (Hu et al., 2006; Lourdin et al., 1997; Qin et al., 2012; Ramiah, 1970). In this chapter, we will review recent accomplishments of thermal analysis of natural fibers and illustrate new methods developed for natural fibers as well as their regenerated biomaterials. Some typical protein and polysaccharide materials, such as silkworm and spider silk fibroins, wool keratins, and cotton celluloses, will be selected as examples to explain how advanced thermal analysis methods are used to study natural fibers, especially protein fibers. We hope these thermal analysis methods can provide scientists a broad view of how polymer thermodynamic theories were used in recent natural fiber studies. This chapter will be organized as follows. We begin with a brief explanation of the natural fiber structures and their self-assembly mechanisms, followed by an overview of thermal analysis on natural fibers and their related materials. We will then focus on different advanced DSC methods and their related applications. This part will be divided into six sections, illustrating how thermal analysis was used to characterize the crystallinity, melting, heat capacity, and crystallization kinetics of natural polymers, as well as how novel thermal analysis methods were developed to model the dynamic biopolymer-bound water systems and biopolymer
metal ions systems.

7.2

Structure of natural fibers

Silks, keratins, collagens, elastins, corn zeins, celluloses, and lignins are common natural fibers (Fig. 7.1), and each of them has distinguished properties. They could be processed into various materials with unique characteristics. Their secondary

Thermal analysis of natural fibers

107

Figure 7.1 Typical raw materials of natural fibers include wool keratin, thai silk, mori silk, and corn zein. Source: Non-copyrighted own figure of Hu.

structures, such as β-sheets in silks (Knight et al., 2000), triple helix in collagens (Brodsky and Ramshaw, 1997), or twofold screw-axis symmetry in cellulose (French and Johnson, 2009), have direct correlation to their properties and applications. Therefore getting a comprehensive and detailed understanding of their hierarchical structures is vital to make a full use of these green and high-performance materials in the future.

7.2.1 Silk Silk proteins are usually produced in the glands of silkworms, spiders, or other insects and are then spun into natural protein fibers. Depending on the specific insect specimen, the structure, chemical composition, and mechanical properties of the specific silks vary significantly. Specifically, silkworm silks and spider silks have been studied for many years and scientists have gained comprehensive knowledge of the structure of the fiber (Gosline et al., 1986; Jiang et al., 2006; Pe´rezRigueiro et al., 1998; Shao and Vollrath, 2002; Xu and Lewis, 1990). For example, silkworm silk fibers have been commercially used as clothing materials for decades (Ricci et al., 2004), and their regenerated materials have been used as biomaterials in biomedical fields (Murphy and Kaplan, 2009). Silkworm silk fibers principally consist of two types of proteins: fibroin and sericin. Fibroin is the main filament of the silk, which accounts for the excellent elasticity of silk fibers due to highly organized β-sheet crystals in the semicrystalline regions of fibroin fibers (Gosline et al., 1999; Santin et al., 1999). Sericins are a kind of sticky hydrophilic and gelatinous proteins, which hold two strands of fibroin fibers together, helping silkworms form the cocoon case that protects their transformation process in nature (Garel et al., 1997). Sericins can be easily removed from fibroin silk fibers by boiling the cocoons in sodium carbonate aqueous solutions. Analysis of spider silk proteins also revealed important information of their structures. Spider silk protein is a self-assembling biopolymer, which consists of numerous but small, organized, and oriented β-sheet crystals (Hinman et al., 2000; Termonia, 1994). It was also found that their protein structure is highly ordered

108

Thermal Analysis of Textiles and Fibers

even in the less crystalline regions. By using Nephila pilipes spider silk as an example, the diameter of silk thread is about 4
5 μm, which consists of a number of silk fibrils with a diameter of 40
80 nm (Du et al., 2006). One silk fibril segment is formed by several β-sheets connected by non-β-sheet structures, and many of these interlinked fibril segments compose one silk fibril along the silk thread direction. Due to this structure, the mechanical properties of spider silks are superior to those of many synthetic materials.

7.2.2 Keratins Keratin is a typical intermediate filament protein existing in epithelia, hair, feathers, nails, horns, hoofs, scales, and wools. Due to its recyclability and unique functions, keratin materials are becoming more and more important in materials science and numerous studies were carried out on various keratins recently. In the epithelial cells, keratins are the most profuse proteins and form the intermediate filament proteins (Wu and Rheinwald, 1981), which generate cytoplasmic network and protect epithelial cells from stress (Barone et al., 2005; Moll et al., 2008). Based on the sequence types, these multigene proteins can be divided into type I (K9
K23; Ha1
Ha8) and type II (K1
K8; Hb1
Hb6). Wool keratins are one of the most useful biopolymers, which have been widely used in wool or leather products as neatening reagents, textile fibers, or additives in animal feeds (Freedberg et al., 2001; Magin et al., 2007). The secondary structures of keratins, mainly, are α-helix in the hair and wool, but β-sheets, in the feather (Gupta and Ramnani, 2006). Depending on the sulfur content, keratins may have high stiffness and inextensibility but could also become very pliable.

7.2.3 Celluloses Polysaccharide is a biopolymer consisting of cellulose, starch, chitin, etc. Because of their chemical reactive groups in the chains, such as hydroxyl, amino, and acetamido groups, polysaccharides are receiving more and more attention for application as additives or adsorbent materials (Samayam et al., 2011). Cellulose is a longchain polysaccharide consisting of thousands of D-glucose units linked by β (1!4)glycosidic bonds. It is one of the most abundant biomasses existing in the plants (Brown, 2004). Cellulose is the major component of cotton (98%) and wood (40%
50%) and is dominant in about one-third of plant tissues. It is estimated that 1011 MT of cellulose are synthesized every year on Earth (Samayam et al., 2011). Cellulose can be divided into cellulose I and cellulose II structures according to its crystalline forms. Cellulose I structure plays the dominant role in cellulosic materials. In cellulose I conformation, chains are bound firmly together to form microfibrils with hydrogen bonds between the multiple hydroxyl groups on the glucose and oxygen atoms. The fibrillar crystallites are partially oriented and can be used as carbon-based feedstock for industrial products such as fuels. Besides, the microfibrils in cellulose can form different shapes, such as giant squares, large rectangular shapes, or thin membrane-like structures, found in many plants. In recent years,

Thermal analysis of natural fibers

109

cellulose nanocrystals and nanofibrils were also developed as renewable materials for various applications (Crini, 2005; Mora´n et al., 2008).

7.2.4 Collagens Collagens are a group of structural proteins (Gelse et al., 2003) accounted for a dominant proportion in the extracellular matrix, with more than 20 types identified. Collagens are usually found in connective tissues or interstitial tissues of parenchymal organs, providing mechanical stability for these tissues. Collagen macromolecules can form a triple helix structure consisting of three polypeptide chains (Kadler et al., 2007; Yokota et al., 2000). In some cases, all three polypeptide chains are identical, but two or even all of them could be different, depending on the collagen types. Besides, the properties of the chains can vary with different amino acids’ sequence of the protein chain. Despite this, all three chains are righthanded helix (Myllyharju and Kivirikko, 2004). Among all kinds of collagens, type I collagen is the most abundant protein in tissues (Hotary et al., 2000; Mizuno et al., 2000). The three helices of type I collagen are hererotrimers consisting of two identical α1(I)-chains and one α2(I)-chain, which provide torsional stiffness, load bearing, and tensile stiffness in the connective tissues.

7.2.5 Other natural fibers There are several other natural fibers that have also been investigated in detail during the past decade. Lignin is a complex aromatic copolymer synthesized from three hydroxycinnamyl alcohol monomers (Guille´n et al., 2005; Zakzeski et al., 2010). It is the second most abundant form of aromatic carbon in the biosphere, existing in the middle layers of cell walls of plants and providing stiffness and strength for the stem. Elastin is an elastic protein found in connective tissues. Elastic fibers in extracellular matrix provide mechanical elasticity to tissues during contracting and stretching. The elastin fibers are formed by cross-linking and assembling tropoelastin monomers together (Uitto, 2008; Yanagisawa et al., 2002). In addition, there are many natural fibers used as green energy materials. With the depletion of fossil energy, more advanced methods were developed to use natural fibers as natural resources for biomass. All these new developments indicated that natural fibers will play a more important role in the energy and material fields in the future.

7.3

Thermal analysis of natural fibers

DSC is a technique widely used to determine thermally induced phase transitions and conformational changes of natural fibers such as silk, cellulose, chitin, and collagen. Since physical transitions and chemical reactions are usually accompanied with heat release or absorption, DSC can measure difference in the heat flow rates

110

Thermal Analysis of Textiles and Fibers

to/from the sample and the reference. The difference in the heat flow rates is proportional to the specific heat capacity of the sample. Therefore DSC can determine the glass transition temperature and the heat capacity increase at the glass transition, the melting and boiling point, the heat of fusion and heat of vaporization, and the crystallization temperature in nonisothermal crystallization experiments. In addition, in isothermal crystallization measurements, DSC can often determine the type of nucleation and crystal growth. Fig. 7.2 shows a typical DSC curve for a biopolymer sample. Fig. 7.3 shows nonisothermal DSC heating curves of various silk samples, including Chinese mulberry (Bombyx mori), Indian Antheraea mylitta (Tussar), Antheraea assama (Muga), and Philosamia ricini (Eri) silkworm cocoons and their degummed fibers. The samples were heated at 2 C/min from 230 C (243K) to 400 C (673K) with temperature regions related to bound water evaporation (Tw), glass transition (Tg), and sample degradation (Td1 and Td2) (Mazzi et al., 2014). There are many advantages of using thermal analysis for studying natural polymers. First, only a very small amount of sample (e.g., 1
5 mg) needs to be used. Second, samples that are not suitable for other types analytical methods can be tested by thermal analysis. For example, samples that cannot be cast onto dissolvable salt plates, or whose surfaces are too rough for attenuated total reflectance measurements of FTIR spectra, can be analyzed by thermal analysis. Second, if the crystallite size of biopolymer materials is very small and their size distribution is very broad, they cannot be readily investigated by X-ray diffraction. In these cases, thermal analysis would be preferential for characterizing these natural polymers. Thus due to the larger sample size, the scattering of the results from thermal

Figure 7.2 An ideal DSC nonisothermal heating curve of a semicrystalline biopolymer material (e.g., silk materials). Unpublished figure by Xue, Hu and Hu.

Thermal analysis of natural fibers

111

Figure 7.3 DSC heating curves of various silk samples, including Chinese mulberry (Bombyx mori), Indian Antheraea mylitta (Tussar), Antheraea assama (Muga), and Philosamia ricini (Eri) silkworm cocoons and degummed fibers. The samples were heated at a rate of 2 C/min. The temperature regions of bound water evaporation (Tw), glass transition (Tg), and sample degradation (Td1 and Td2) are indicated in the figure. Source: From Mazzi, S., Zulker, E., Buchicchio, J., Anderson, B., Hu, X., 2014. Comparative thermal analysis of Eri, Mori, Muga, and Tussar silk cocoons and fibroin fibers. J. Therm. Anal. Calorim., 116(3), 1337
1343, there Fig. 2. Fig. 7.4; Hu, X., Kaplan, D., Cebe, P., 2006. Determining beta-sheet crystallinity in fibrous proteins by thermal analysis and infrared spectroscopy. Macromolecules, 39(18), 6161
6170, Fig. 5b. Reprinted with permission from Springer International Publishing AG. This experiment is performed using a TA Instruments Q100 DSC.

analysis measurements is considerably smaller than those collected from a single point scan of various spectroscopes (e.g., FTIR and Raman spectroscopes).

7.3.1 Measurement of crystallinity of natural fibers by differential scanning calorimetry The crystallinity of natural fibers has a great impact on their mechanical properties (Bergander and Salme´n, 2002; Park et al., 1999). Therefore the determination of the crystallinity of natural fibers is an important part of the biopolymer structural
functional analysis. These crystals are usually stacked chains aggregating with ordered structures and are assembled in the natural polymers (Hu et al., 2009b). The physical and chemical properties of natural fibers will vary with varying crystallinity. Many analytical methods were developed for measuring the crystallinity of natural polymers, including the X-ray scattering and various spectrum deconvolution methods.

112

Thermal Analysis of Textiles and Fibers

7.3.1.1 Glass transition Thermal analysis provided a new investigative method for the study of natural polymer structure and/or crystallinity, which is based on the measurement of specific heat capacity increment at the glass transition (Menczel and Jaffe, 2007; Menczel and Wunderlich, 1981, 1986; Turi, 2012). This method is well known in the study of synthetic semicrystalline polymers for determining the “mobile amorphous fraction” of the semicrystalline polymers, which can exhibit its glass transition relaxation process during heating (Arnoult et al., 2007). An example in recent studies of natural biomaterials is the β-pleated sheet crystal in fibrous proteins (Hu et al., 2006; Qin et al., 2012). These hydrophobic β-sheet crystals constitute the nucleation site for protein folding in the protein fibers and control their unfolding process by the relaxation of these highly thermally resistant structures in the later stage of the unfolding pathway. Determination of the heat capacity increases at the glass transition. Solid curve: DSC heating curve of a semicrystalline biopolymer; dashed curve: DSC heating curve of a noncrystalline biopolymer. The solid straight lines represent the temperature dependence of the specific heat capacity of the crystalline [Cp(T)solid] and liquid phases [Cp(T)liquid]. Cp(T)SC is the specific heat capacity of a semicrystalline polymer between the glass transition and the melting point. This is smaller than Cp(T)liquid due to the presence of the crystalline phase and the rigid amorphous phase. These specific heat capacity lines are extrapolated to the glass transition temperature and the difference between them is called “the heat capacity increase at the glass transition” (ΔCp). ΔCp100% is the heat capacity increase at the glass transition for a 100% amorphous (i.e., zero crystallinity ΔCp100% is denoted as ΔCp0 in Fig. 7.4.) polymer, while ΔCpSC is the heat capacity increase at the glass transition for a semicrystalline polymer (smaller than ΔCp100% due to the crystallinity and the presence of the possible rigid amorphous phase). Fig. 7.4 indicates how to determine the heat capacity increase at the glass transition for an amorphous and semicrystalline polymer. The amorphous phase is a glassy state below the glass transition temperature. When the temperature rises to temperatures above the glass transition temperature, the amorphous phase devitrifies, that is, the glass becomes melt phase. This process in the DSC curve is observed as a heat capacity jump. Meanwhile, the crystalline fraction does not change up to the melting point where the crystals become unorderly melt. Therefore the amorphous fraction Φam can be determined by measuring the heat capacity increase at the glass transition (Tg): Φam 5

ΔCpSC ðTg Þ ΔCp100% ðTg Þ

(7.1)

where Φam is the heat capacity increase at Tg, a semicrystalline biopolymer. And ΔCp100%(Tg) is the heat capacity increase at Tg for a fully amorphous biopolymer, which can sometimes be prepared by casting very dilute biopolymer aqueous solution into thin film. FTIR and X-ray diffraction methods can also be used to confirm and verify the completely amorphous state of the biopolymer samples and thus

Thermal analysis of natural fibers

113

Figure 7.4 The approach to determining the heat capacity increment at the glass transition temperature, Tg. Dashed curve: less ordered, noncrystalline natural polymers; light solid curve: crystalline natural polymers containing crystals; heavy solid line: the tangents to the heat capacity curve at low and high temperature above Tg. The solid-state heat capacity, Cp(T)solid, the liquid heat capacity, Cp(T)liquid (from the noncrystalline sample), and the semicrystalline heat capacity, Cp(T)SC (from crystallized sample), can be derived from the tangents. The increment is determined from the difference between the extrapolated tangents of the heat capacity at the Tg. Source: From Hu, X., Kaplan, D., Cebe, P., 2006. Determining beta-sheet crystallinity in fibrous proteins by thermal analysis and infrared spectroscopy. Macromolecules, 39(18), 6161
6170, Figure 5b. Reprinted with permission from American Chemical Society. The experiment is performed using a TA Instruments 2920 DSC.

based on the common two-phase models, one can directly obtain the fraction of the crystalline phase using the following equation: "

ΔCpSC ðTg Þ Φc 5 1 2 Φam 5 1 2 ΔCp100% ðTg Þ

# (7.2)

In addition (Lau et al., 1984; Grebowicz et al., 1984; Suzuki et al., 1985), all semicrystalline polymers (including biopolymers) can also be modeled as consisting of three fractions: the mobile fraction, ΦM, the crystalline fraction, Φc, and the immobilized noncrystalline fraction (rigid amorphous fraction) (Cheng et al., 1986; Huo et al., 1993; Lu and Cebe, 1996; Menczel and Jaffe, 2007; Menczel and Wunderlich, 1981). The immobilized amorphous fraction ΦIMM, called the “rigid amorphous fraction” in semicrystalline synthetic polymers, includes any immobilized noncrystalline portions

114

Thermal Analysis of Textiles and Fibers

of the natural polymer that remain immobile above the glass transition temperature of the mobile fraction (i.e., the crystal
amorphous interface and the tight tie molecular fraction). From here, it follows that the existence of the rigid amorphous phase is not absolutely necessary in a semicrystalline polymer, although it is a frequent phenomenon. If the surface of the crystallites is ideal, for example, regular chain folding takes place throughout the crystalline surfaces, the rigid amorphous phase may be missing. Thus if we can obtain the crystallinity of biopolymer fibers from other analytical techniques, such as FTIR or X-ray scattering, or the melting peak from the DSC heating experiments, we can also model the three phases and their fractions. In the study of the B. mori silk fibroin proteins, a two-phase model was found from the various measurements. This heat capacity method can be broadly utilized to understand structures of other natural polymers as well.

7.3.1.2 Blending method Gomba´s et al. (2002) reported that another DSC method can be used to test the crystallinity of crystalline/amorphous powder mixtures such as α-lactose monohydrate, which is well known for its existence in crystalline and amorphous states. The authors first prepared amorphous and 100% crystalline lactose. Then, the amorphous and crystalline lactose were blended to prepare standard samples with crystallinity values of 0, 5, 10, 20, 30, 40, 50, 60, 70, 80, 90, 95, and 100 mass%. By running the DSC experiments of the 100% amorphous and 100% crystalline samples, they found that an exothermic peak (II) at 167 C, which indicated the transformation of amorphous state to crystalline state, and an endothermic peak (I) at 144 C, which indicated the loss of crystalline water, is typical for the crystalline lactose. Therefore as shown in Fig. 7.5A, with the content of the crystalline

Figure 7.5 (A) DSC curve of a lactose sample with 50% crystallinity. (B) Relation between actual and predicted crystallinity of lactose samples. DSC, Differential scanning calorimetry. Source: From Gomba´s, A., Szabo´-Re´ve´sz, P., Kata, M., Regdon, G., & Er˝os, I., 2002. Quantitative determination of crystallinity of α-lactose monohydrate by DSC. J. Therm. Anal. Calorim., 68(2), 503
510, Figs. 5 and 6. Reprinted with permission from Akade´miai Kiado´ Zrt. The experiment is performed using a Mettler-Toledo DSC 821.

Thermal analysis of natural fibers

115

component increasing, the magnitude of the exothermic peak will decrease, but the size of the endothermic peak will increase. According to this experimental method, researchers can use two parameters and the transition energy values of amorphous part in the samples to evaluate the crystalline degree of a semicrystalline biopolymer sample quantitatively in the future (Young et al., 2000). For example, Gomba´s et al. (2002), in this study, created an equation using a regression method for predicting the crystallinity of lactose materials: y 5 87.241 2 0.825x, where y is the predicted crystallinity degree (%), x is the transition energy (J/g), 87.241 is the intercept, and 0.825 is the slope. This method is useful for physically mixed polymer composites with crystallines and noncrystallines, whose amorphous content is at least 20%, while its accuracy can be confirmed by other methods such as X-ray diffraction. Although this method might not be useful for a real semicrystalline natural polymer, it provided an easy way to get the crystallinity of lactose mixtures by only the transition energy value of its amorphous part.

7.3.2 Thermal analysis characterization of the melting of natural fibers Natural fibers have extraordinary physical properties due to the interchain hydrogen bonds formed in the fibers. For example, β-pleated sheet crystals found in fibrous proteins are one of the most stable crystal structures in natural fibers (Holland et al., 2012; Omenetto and Kaplan, 2010), which were too stable to be melted by a regular DSC through input of heat energy (Cebe et al., 2013). This also accounts for the difficult cure of Alzheimer’s disease. However, it was recently revealed that using the fast-scanning chip calorimetry at a heating rate of 2000 C/s, we can achieve the reversible thermal melting of β-pleated sheet crystals above 530K (B260 C) (Cebe et al., 2013). This finding demonstrated the feasibility of thermal melting of any stable crystal in natural fibers. By changing the crystal structure in natural fibers with heat treatment, one can easily control the properties of natural fibers in the future. Cebe et al. used silk fibers of domesticated silkworm B. mori as exemplary structural protein for studying the melting of two-dimensional β-sheet crystals (Fig. 7.6A) (Altman et al., 2003; McGrath and Kaplan, 1997; Zheng et al., 2010). They found that the silk melting transition is obscured by decomposition when it is heated at 1
20 C/min using regular DSC. The thermal decomposition began at 470K. Hence, the intrachain bonds begin to break well before the melting transition is over (Fig. 7.6B) (McGrath and Kaplan, 1997; Vepari and Kaplan, 2007). However, the F-DSC can heat or cool a sample at a rate of several thousand degree Celsius per second (Mathot et al., 2011; Van Herwaarden et al., 2011). Due to the extremely fast heating and cooling, the exposure time of samples for high temperatures is greatly reduced and the whole scanning circle takes only B0.3 seconds. Therefore this technology provides an advantage to observe the melting of silk crystals. When a silk fibroin sample, containing β-pleated sheet crystals, was heated at 2000K/s in nitrogen atmosphere, its Tg (474K 5 201 C, at this heating rate) and Tm

116

Thermal Analysis of Textiles and Fibers

Figure 7.6 (A) During fast heating, β-pleated sheet crystals can be melted in 0.002 s. (B) Bombyx mori silk film is heated at 2K/min in DSC. (C) B. mori silk film is heated at 2000K/ s using fast-scanning chip calorimetry. In both heatings, the heating rate is 2000K/s. (D) Heat capacity versus temperature, at 2000K/s, for a crystallized silk fibroin film and a bundle of native degummed cocoon fibers. Down and up arrows represent Tg and the beginning of melting, respectively. The horizontal double arrow shows the region of fiber melting and shape change to droplet morphology. Source: From Cebe, P., Hu, X., Kaplan, D. L., Zhuravlev, E., Wurm, A., Arbeiter, D., Schick, C., 2013. Beating the heat-fast scanning melts silk beta sheet crystals. Sci. Rep. 3, 1130, Figs. 1c, 2a, 2c, and 3c. Reprinted with permission from Nature Publishing Group. The experiment is performed using a self-developed Fast Scanning Calorimetry (FSC) with a thin film chip sensor XI-399 of Xensor Integration (The Netherlands).

(endothermic peak around 530K 5 257 C) were, for the first time, displayed together (Fig. 7.6C). A similar experiment was also conducted on degummed native cocoon silk fiber (Fig. 7.6D). The described fast-scanning chip calorimetric technique has many advantages: it is used with sample masses as little as 10 ng. The sample is put directly on the 500 nm thin SiNx membrane and all fast-scanning experiments are conducted in nitrogen atmosphere without the disturbing effect of water vapor. The differential power supplies heat necessary for keeping the sample and the reference temperature at the same level. Therefore the output signal of the instrument is proportional to the specific heat capacity. This thermal analysis method can be applied to other natural fibers, where their structures have a great influence of its physical and

Thermal analysis of natural fibers

117

chemical properties. This technique can also help us further understand the reorganization of other time-dependent processes in natural fibers.

7.3.3 Thermal analysis measurement of the specific heat capacity of natural fibers The heat capacity measurements of the biopolymers are very important in the modern biotech research and industry. For example, the polymerase chain reaction (PCR) technique has become the major achievement in creating DNA molecules for the industry and research products (Holland et al., 1991; Mullis et al., 1986, 1994). The DNA structure changes in the PCR process are simply controlled by adjusting its temperatures with the help of salts and enzymes in the solution (Saiki et al., 1985). The selected temperatures, for example, the melting temperature of DNA, were directly predicted by the numbers and heat capacities of the four fundamental DNA bases—A, T, G, C in a certain salt solution. Similarly, the effect of heat capacity of the biopolymers, especially the protein molecules, is important when studying protein-based natural fibers. The ambient temperature and heat capacity of the protein molecules can affect their final biofunctions in the living body. However, for the most complicated and cherished protein families, separated heat capacity study based on each protein structure will be time-consuming. This problem leads the researchers toward exploring practical ways for predicting the heat capacity properties of large, complicated protein materials.

7.3.3.1 Theoretical prediction Since the amino acid constituents of the proteins are few in number, and for many physical properties (such as vibrational heat capacities) a simple additivity rule may be applied to predict the behavior of the macromolecule. By using a linear combination, the bulk property of the large macromolecule, P, can be predicted from its P small constituent properties, p, through: P 5 iNiPi, where Ni is the number of units characterized by each property Pi. Using this method, Pyda et al. (Huang et al., 2011; Pyda et al., 2008) predicted the vibrational motions of heat capacities of some natural protein materials based on their unique amino acid components from 0K to the glass transition region of the polymer. First, the total solid vibrational heat capacity of a protein, Cpprotein , is constructed as a sum of products of the vibrational heat capacity of the individual poly(amino acid) residues, Cp(i), with each of their numbers found in the complete sequence of amino acids in the protein macromolecule: P Cpprotein 5 i Ni Cp ðiÞ 5 ðNGly Cp ðGlyÞ 1 NAla Cp ðAlaÞ 1 ? 1 NMeth Cp ðMetÞÞ

(7.3)

where Ni is the total number of each type of amino acid and can be calculated from:

118

Thermal Analysis of Textiles and Fibers

N i 5 Xi

Mw ðproteinÞ Mw ðiÞ

(7.4)

where Xi is the molar fraction of each type of amino acid unit in this protein; Mw(protein) is the total molecular weight mass of the protein; and Mw(i) is the molecular weight mass of each kind of amino acid according to the protein amino acids analysis (e.g., i 5 Gly, Ala, Ser, . . ., Met). It needs to be noted that Mw(i) is the molecular weight mass of the repeating unit of poly(amino acid)s rather than the amino acid’s small molecule. Heat capacity, Cp(i), is the vibrational heat capacity of the repeating unit of the ith kind of poly(amino acid) in the solid state, which is collected in a polymer heat capacity database and the numbers, Ni, and compositions of the amino acids in the protein can be obtained from worldwide protein sequence data banks (Bairoch and Apweiler, 2000; Boeckmann et al., 2003). Fig. 7.7 shows a comparison of calculated vibrational heat capacity of protein, Cp(vibration)protein (solid curve) with the experimental heat capacity of noncrystalline protein (open circles) using B. mori silk fibroin (Fig. 7.7A) and recombinant spider silk materials (Fig. 7.7B) as examples. The vibrational heat capacity of protein, Cp(vibration)protein, was evaluated using the vibrational motion spectra of individual amino acids in silk proteins. This vibrational heat capacity was constructed according to Eq. (7.3) using the vibrational Cp of different poly(amino acid)s from a data bank. Knowing the numbers Ni of each type of amino acid in the protein and their corresponding vibrational heat capacity as function of temperatures, the (B) CpSilk(liquid)

1200 Cp(experimental)

800

400

0

CpGly+Ala+Ser(vib)

CpSilk(vibration)

CpGly(vib)

CpAla(vib) CpSer(vib)

0

100

200

300

400

Temperature (K)

500

Rev. heat capacity (kJ/K/mol)

Heat capacity (kJ/K/mol)

(A)

600

CpHAB3(liquid)

30 25 20

CpHAB3(experimental)

Tg 447.1K

15 10

CpHAB3(vibration)

5 0 0

CpA(vibration)

73

173

CpB(vibration)

273 373 Temperature (K)

CpH(vibration)

473

573

Figure 7.7 Comparison between measured heat capacity (open circles) and calculated vibrational heat capacity (thick solid curve) as functions of temperature for (A) Bombyx mori silkworm silks and (B) recombinant spider silk
like block copolymers. Vibrational heat capacities of the amino acids, glycine (Gly), alanine (Ala), and serine (Ser), and the sum of them, are indicated. Source: From Pyda, M., Hu, X., Cebe, P., 2008. Heat capacity of silk fibroin based on the vibrational motion of poly(amino acid)s in the presence and absence of water. Macromolecules, 41(13), 4786
4793, Fig. 2; Huang, W., Krishnaji, S., Hu, X., Kaplan, D., Cebe, P., 2011. Heat capacity of spider silk-like block copolymers. Macromolecules, 44(13), 5299
5309, Fig. 4a. Reprinted with the permission from American Chemical Society. Both are performed using a TA Instruments Q100 DSC.

Thermal analysis of natural fibers

119

Figure 7.8 Heat capacity of cellulose I (cotton) and water systems. Source: From Hatakeyama, T., Nakamura, K., Hatakeyama, H., 2000. Vaporization of bound water associated with cellulose fibres. Thermochim. Acta, 352, 233
239, Fig. 9. Reprinted with permission from Elsevier Science B.V. The experiment is performed using a Seiko DSC 200.

vibrational heat capacity baseline of protein, Cp(vibration)protein, can be calculated over the whole temperature range of 0.1K
600K. Good agreement between the experimental data [particularly when the contribution of the rotational heat capacity in proteins (only due to the presence of certain pendant groups) is very small] and the calculated vibrational heat capacity protein materials are shown below its glass transition temperature, Tg of 451.15K (5178 C). This heat capacity prediction model could be widely useful in the heat capacity studies of other proteins or natural polymers.

7.3.3.2 Experimental measurement Besides the theoretical calculations described earlier, many experimental researches were also developed to calculate heat capacity of a complicated natural material system. For example, Hatakeyama et al. (2000) measured the specific heat capacity of natural cellulose I-water system by DSC. The results are shown in Fig. 7.8, where Wc stands for water content (mass of water/mass of dry cellulose).

7.3.4 Thermal analysis characterization of the crystallization kinetics of natural polymers Polymer crystallization is a phase transition, which can be characterized by thermal analysis. DSC can be used to determine the kinetics of this transition and provide

120

Thermal Analysis of Textiles and Fibers

us information about how polymer crystallization takes place dynamically (crystallization speed, size, dimension, and distribution, etc.) (Piorkowska et al., 2006; Schultz, 2001; Wunderlich, 1976). This kinetics process can be described by wellknown theories such as the Avrami equation (Schick, 2009; Wunderlich, 1976) The application of the Avrami equation can provide such important data as the type of ˇ ´ k and Berggren, 1971; Uhlmann, nucleation and geometry of growing entities (Sesta 1972). Proteins and polysaccharides are essentially a group of natural or biosynthesized linear polymers (McGrath and Kaplan, 1997). These materials can stay amorphous or be transferred into ordered structures such as three-dimensional crystals; thus the macroscopic properties of these materials can sometimes be changed significantly by changing the ordering of macromolecules. However, it is usually very difficult to calculate the crystallization kinetic processes of biopolymer crystals, such as β-sheet crystals, directly in protein. One of the reasons is that these crystals are too small to be observed (e.g., the size of the β-sheet crystal regions is only B10 nm) (Drummy et al., 2007; Jin and Kaplan, 2003; McGrath and Kaplan, 1997). Therefore thermal analysis methods could be very useful for observing and calculating the kinetics of biopolymer crystal growth. For example, the isothermal crystallization kinetics of silk protein was investigated using DSC performed around the crystallization temperatures. Fig. 7.9A shows isothermal crystallization curves of silk fibroin at four temperatures between 192 C and 202 C. As it can be seen in Fig. 7.9A, the crystallization process lasts up to time values three to four times greater than the peak time of the exothermic (A)

(B) 1 202ºC Relative crystallinity

Heat flow (a.u.)

202ºC

199ºC

195ºC

199ºC 195ºC

192ºC

0.8 0.6 0.4 0.2

192ºC

0

50 100 Time (min)

150

0 0

50

100 150 Time (min)

200

Figure 7.9 (A) DSC curves of isothermal crystallization of silk fibroin at four different temperatures: 192 C, 195 C, 199 C, and 202 C. (B) Relative crystallinity (Xrel) versus crystallization time for silk fibroin in isothermal crystallization at the temperatures mentioned earlier. Source: From Hu, X., Lu, Q., Kaplan, D. L., Cebe, P., 2009b. Microphase separation controlled β-sheet crystallization kinetics in fibrous proteins. Macromolecules, 42(6), 2079
2087, Figs. 2 and 3. Reprinted with the permission from American Chemical Society. The experiment is performed using a TA Instruments Q100 DSC.

Thermal analysis of natural fibers

121

crystallization peak. If we denote the crystallinity at infinite times as XN, then Xt/XN will be called relative crystallinity at any point of time. The relative crystallinity (Xrel) at any point of time can be obtained from the following equation (Choi and Kwak, 2007; Xu et al., 2002): Xxrel ðtÞ 5

Ð t   H ðt Þ 0 dt dH ðtÞ=dt   Ð 5 N H ðNÞ 0 dt dH ðtÞ=dt

(7.5)

where H(t) is the heat evolved until time t, H(N) is the total heat of crystallization, and t is time. Fig. 7.9B shows the time dependence of the relative crystallinity as measured by DSC. All curves have sigmoidal shapes. The original Avrami equation has the following form (Avrami, 1939; Miles et al., 1985): XðTÞ 5 1 2 expð2 KT n Þ

(7.6a)

where Xrel is the relative crystallinity of the crystallizing sample at different time t, K is the rate constant of crystallization, and n is the Avrami exponent, containing information related to the nucleation mechanism and geometry of the growing crystallites (Schulz and Wunderlich, 1976). K and n can be determined from the graph if log[ 2 ln(1 2 Xrel)] is plotted against log(t). This is the original form of the Avrami equation, which can be used for lowmolecular-mass substances if the crystallization starts as soon as the sample reaches the crystallization temperature. Later, the Avrami equation was modified for polymer crystallization experiments. First, the relative crystallization (Xrel) was introduced because polymers never have 100% crystallinity after melt crystallization experiments. Second, the induction time of crystallization was inserted into the equation. This question is still debated in the literature. Several authors (see, e.g., Varga et al., 1984) indicated some proof that the Avrami equation should be applied to the crystallization process starting with time when the crystallization process starts. The time between the start of the crystallization process and zero time (i.e., when the sample temperature reaches the crystallization temperature) is called the induction time of crystallization, τ i. With these modifications the Avrami equation for polymer melt crystallization studies takes the following form: X rel 5

XðtÞ 5 1 2 exp½ 2Kðt2τ i Þn  XðNÞ

(7.6b)

Very often, when log[ 2 ln(1 2 Xrel)] is plotted against log(t 2 τ i) in polymer crystallization studies, the lines are more linear (i.e., the scattering of the points is smaller) and the obtained Avrami exponent can be better used to explain the crystallization process. Usually, the linear region of the curves is taken to determine the crystallization rate constant K and Avrami exponent n. Combined with the kinetics data obtained

122

Thermal Analysis of Textiles and Fibers

from real-time Wide-angle X-ray scattering (WAXS) and real-time FTIR methods, we could obtain a unique Avrami exponent of biopolymers (e.g., for this mori silk, nB2 indicating two-dimensional crystal growth for heterogeneous nucleation). A crystallization kinetic model can be further developed to explain the crystallization mechanism in this natural polymer. For this example, the unique, repeating amino acid sequence arrangement in the natural silk protein is very similar to many A
B semicrystalline block copolymers, in which crystallizable regions A are surrounded with noncrystallizable regions B. This peculiarity may be responsible for causing two-dimensional crystallization of silk proteins, resulting in an Avrami exponent of nB2 (heterogeneous nucleation). As crystallizable
noncrystallizable regions are very common in biopolymers, this thermal analysis method is applicable to the crystallization of many natural fibers and provide scientists a new tool for understanding the chain folding process in natural polymers.

7.3.5 Thermal analysis study of the natural polymer
bound water systems One of the major challenges in polymer science is the water
polymer interaction. Most natural proteins or polysaccharides contain bound water at the chain folds on the crystallite surfaces (Hu et al., 2008; Rajkhowa et al., 2012). Water molecules affect not only the stability of the biopolymer three-dimensional structure but also the process of their phase transition and crystallization (Qin et al., 2012). Generally, hydration results from different interactions between water molecules and specific functional groups of the natural polymers (Lee and Ha, 1999; Slade and Levine, 1995). Bound water molecules can act as a plasticizer and become incorporated into the biopolymer chains by connecting up to four hydrogen bonds to increase flexibility, extensibility, and workability of the natural fibers. Therefore bound water can expand the accessible conformational space of the biopolymer by decreasing its inter- and intramolecular friction and effective barrier heights for conformational change. This will result in a new (lower temperature) “glass transition” for the whole biopolymer
water coupling system. With the water molecules transferring across the protein molecular layers, a structure transformation in proteins could be induced. This dynamic protein-bound water transition can be observed in detail and analyzed by thermal analysis methods (Wortmann et al., 2006). Using silk fibroin protein as an example, Fig. 7.10A shows DSC scans of silk fibroin proteins at various heating rates. The samples from Fig. 7.10A (a)
(d) are silk fibroin films containing about 5.5 wt.% bound water (Hu et al., 2007). The sample in Fig. 7.10A (e) was pure silk fibroin without bound water as a reference sample (Hu et al., 2006, 2007), having a glass transition temperature of about Tg(2) 5 178 C and a nonisothermal crystallization peak(2) at around 213 C. After the crystallization peak, the protein started to degrade with endotherm peak at around 260 C. There were no endothermic or exothermal events found below the glass transition temperature Tg(2) for the pure silk fibroin sample. However, the silk fibroin samples containing bound water

Thermal analysis of natural fibers

123

Figure 7.10 (A) Silk fibroin films are heated from 265 C to 280 C with the rates 20, 10, 5, and 2K/min. Silk fibroin films from (a) to (d) contain 5.5 wt.% bound water. The sample (e) is a pure silk fibroin after water removal. (B) Curve A silk sample heated at a rate of 5K/ min, modulated with an amplitude of 0.796K over a period of 1 min; Curve B silk sample heated at a rate of 2K/min, modulated with amplitude of 0.318K over a period of 1 min. Dashed lines of Curves A0 and B0 reflect the mass-corrected heat capacity of samples (A) and (B), respectively. Source: From (A) and (B) Hu, X., Kaplan, D., Cebe, P., 2008. Dynamic protein
water relationships during β-sheet formation. Macromolecules, 41(11), 3939
3948, Figs. 1 and 2. Reprinted with permission from American Chemical Society. The experiment is performed using a TA Instruments Q100 DSC; (C) Wortmann, F.J., Stapels, M., Elliott, R., Chandra, L., 2006. The effect of water on the glass transition of human hair. Biopolymers, 81(5), 371
375, Fig. 1. Reprinted with the permission from Wiley, Inc. The experiment is performed using a Perkin Elmer DSC-7 instrument.

in samples of Fig. 7.10A (a)
(d) exhibited different thermal properties when compared with the dried sample Fig. 7.10A (e). A new additional glass transition, Tg(1), was found for these samples at around 80 C and a series of low temperature exotherm peak(1) was also formed. These initial glass transitions indicated that the bound water has strongly affected the protein structures before it reached its pure glass transition at 178 C.

124

Thermal Analysis of Textiles and Fibers

Fig. 7.10B shows the reversing heat capacity of the silk-bound water samples measured by modulated temperature DSC (MT-DSC). Sample (A) was heated at a rate of 5 C/min, modulated with amplitude 0.796 C over a period of 1 minute; Sample (B) was heated at a rate of 2 C/min, modulated with amplitude 0.318 C over a period of 1 minute. In the temperature region between 80 C and 160 C, a clear glass transition, Tg(1), was formed in both samples (A) and (B) in Fig. 7.10B, indicating silk may first form a new structure during this low temperature region using water as a plasticizer. All of the bound water molecules were finally evaporated around 155 C. This quantitative TMDSC study gives us a clearer picture about what happens when the bound water plasticizer induces a new glass transition in the silk protein
water system and how this structure transforms to the dry solid state of silk with the loss of the water molecules. This “double glass transitions” phenomena was also found in other biopolymer-bound water systems (e.g., a human hair
bound water system has a low glass transition temperature around 40 C, shown in Fig. 7.10C, but the glass transition of pure hair keratin is above 150 C). These studies give us a useful tool in learning how to study crystallization mechanism of other natural polymer
bound water structures in the future.

7.4

Thermal analysis study of the natural polymer
metal ions systems

Proteins can easily bind with metal ions, leading to unique biological properties (Branden, 1999). Serving as the potential ligands for a metallic cation, amino acid residues can control the overall structural stability of proteins in different living circumstances. Many analytical methods, for example, nuclear magnetic resonance and X-ray scattering, have been used to study protein
metal ion interactions (Rose, 1971). With these methods, researchers can obtain the precise information about the metal binding sites in the protein chains as well as the bound strength or the number of metal ions bound to protein molecules. Advanced thermal analysis, as a novel method for protein studies, can also be used to investigate the protein
metal ions systems (Hu et al., 2009a). Since different metal ions have different functions, a selected protein
metal ion structure may have very unique thermal properties. A simple case is the plasticizing and antiplasticizing behavior of metal ions in the proteins. Fig. 7.11A shows the specific heat capacity curve of silk protein-KCl (1 wt.%) system tested by MT-DSC (solid curve) around the temperature of its glass transition. Quasi-isothermal MT-DSC curves are also shown (open and filled circles) in order to get the liquid and solid-state heat capacity values near the glass transition temperature. The measured Tg of the silk-KCl (1.0%) system from MT-DSC curves is 171 C and is a little less than that of pure silk fibroin protein (178 C). For the purpose of easy comparison, the reversing specific heat capacity of pure silk fibroin from a typical MT-DSC heating scan is plotted as a dashed line (Hu et al., 2007, 2008). A theoretical model based on:

Thermal analysis of natural fibers

125

Figure 7.11 Specific heat capacity curves of (A) silk-KCl (1.0 wt.%) and (B) silk-CaCl2 (2.0 wt.%) systems from MT-DSC scans (solid curve) and the quasi-isothermal MT-DSC curves (open and filled circles) in their glass transition regions. Source: From Hu, X., Kaplan, D., Cebe, P., 2009a. Thermal analysis of protein
metallic ion systems. J. Therm. Anal. Calorim., 96(3), 827
834, Figs. 3 and 4. Reprinted with permission from Springer International Publishing AG. Both are performed using a TA Instruments Q100 DSC.

Cpsolid ðprotein-ionÞ 5 XproteinUCpsolid ðproteinÞ 1 XionUCpglassy ðionÞ

(7.7a)

Cpliquid ðprotein-ionÞ 5 XproteinUCpliquid ðproteinÞ 1 XionUCpglassy ðionÞ

(7.7b)

was built with the purpose of studying the heat capacity of the silk protein
metal ions system. Usually, Cp(ion)glassy are close to Cp(ion)liquid, because the metal ion salts with no bound water inside have no melting before the degradation temperature of most proteins. As shown in Fig. 7.11A, the solid and liquid heat capacity baselines of silk protein-KCl system (solid line) were calculated using Eqs. (7.7a) and (7.7b) for the total solid specific heat capacity, Cp(protein-ion)solid fits the DSC curve perfectly in the temperature region below Tg, which means that mostly vibrational motion of silk and ion salt contribute to the heat capacity in this temperature region. As for the total liquid specific heat capacity, Cp(protein-ion)liquid, there is also a good agreement existing if we just use Cp(ion)glassy as Cp(ion)liquid in Eq. (7.7b). [The rotational motion for metal ions can be neglected because they are more-or-less symmetric. And if the ions are bound to the protein chain, there should be no translational motion. So the only motion is vibration, which makes Cp(ion)glassy close to Cp(ion)liquid. Hence, the downshift of the glass transition is the only clear difference between the pure silk fibroin and the silk protein-KCl system. Hence, it means that KCl plays a role of plasticizer in the silk protein structure and causes a decrease in the glass transition of the total system. KCl is a plasticizer because it separates the macromolecular segments from each other, thus leading to

126

Thermal Analysis of Textiles and Fibers

a decreased interaction between them, and this is plasticization. On the contrary, the question is whether one KCl molecule will interact with only one ionic or partial dipole site of the protein. If yes, it is a plasticizer. If it may interact with several groups (or two groups on two macromolecular segments), it can magnify the interaction between the segments and may play the role of physical crosslinks, in which case, Tg should go up. However, other metal ions may have a different effect in the silk protein system. Fig. 7.11B shows the specific heat capacity curve of silk protein-CaCl2(2 wt.%) system from the MT-DSC scan (solid curve) around its Tg. The open and filled circles from quasi-isothermal MT-DSC are used for comparison. From the low temperature to the onset temperature of the glass transition, the heat capacity curves in the solid state match the baseline very well. In the solid state, the quasi-isothermal MT-DSC curve has a sharper slope when compared with the MT-DSC curve, but finally they begin to overlap at 230 C. The calculated heat capacity baselines of silk proteinCaCl2 system (solid line) in both the states are also drawn based on Eqs. (7.7a) and (7.7b) and MT-DSC heat capacity curves of pure silk fibroin (dashed curve) and the calculated solid and liquid state heat capacity baselines of the pure silk protein (dotted lines) are also shown similar to those in Fig. 7.11A. It is clear that the Tg of silk-CaCl2 system is higher than the Tg of pure silk fibroin proteins reaching 197 C. However, the DSC curves fit both solid and liquid silk-CaCl2 noncrystalline baselines very well, indicating zero crystal formation in the silk proteins, meaning that CaCl2 plays a role of an antiplasticizer in the silk protein structure, resulting in an increase in the glass transition of the total system. Compared with the effect of KCl salts, CaCl2 salts could make the structure of silk protein change dramatically in agreement with previous researches (Kaplan, 1994; Tsuda et al., 2002; Zhou et al., 2005a,b). It is of no doubt that structure analysis and microscopy observation can help us get further understanding of ions’ effects on the protein structures. However, thermal analysis is a novel method, which can reveal the impact of the metal ions on the thermal properties of proteins. It is a new pathway in studying protein
metallic ion interactions and has great potential to be applied to other natural fiber structural studies in the future.

7.5

Conclusion

Although many advanced thermal analysis techniques have been developed for the studies of synthetic polymers, few of them were extended to natural fibers, especially for protein and polysaccharide materials. However, with the fast development of the biomedical science and engineering and the fast growth of biomimicry studies, the need for developing modern thermal analysis techniques for these materials became urgent. Advanced thermal analysis techniques such as TGA, TMA, DMA, DSC, F-DSC, and MT-DSC not only provide thermal information on phase transitions and structural changes of the natural fibers but also offer useful information

Thermal analysis of natural fibers

127

on the phase fractions (or crystallinity), heat capacity, and crystallization kinetics of the studied materials. The impact of critical small molecules, such as bound water and metal ions in the natural polymers, can also be studied using these modern thermal analysis techniques. These techniques are also potential candidates for being used as new methods for the dynamic studies of composite biopolymer materials in the future. In this chapter, we broadly reviewed recent developments of novel thermal analysis techniques/methods for characterization of natural fibers and related regenerated materials from these fibers. Several materials, such as silk proteins and celluloses, were used as examples to demonstrate how these different advanced thermal analysis methods can help in understanding the structures and transitions of natural biopolymers. These results and recent ongoing studies on other protein and polysaccharide materials also suggest a possible wide application of these methods in the field of natural or green materials, which would benefit the whole modern thermal analysis society and the biomaterials society in the future.

References Altman, G.H., Diaz, F., Jakuba, C., Calabro, T., Horan, R.L., Chen, J., et al., 2003. Silkbased biomaterials. Biomaterials 24 (3), 401
416. Arnoult, M., Dargent, E., Mano, J., 2007. Mobile amorphous phase fragility in semicrystalline polymers: comparison of PET and PLLA. Polymer (Guildf). 48 (4), 1012
1019. Avrami, M., 1939. Kinetics of phase change. I General theory. J. Chem. Phys. 7 (12), 1103
1112. Bairoch, A., Apweiler, R., 2000. The SWISS-PROT protein sequence database and its supplement TrEMBL in 2000. Nucleic Acids Res. 28 (1), 45
48. Barone, J.R., Schmidt, W.F., Liebner, C.F., 2005. Thermally processed keratin films. J. Appl. Polym. Sci. 97 (4), 1644
1651. Bergander, A., Salme´n, L., 2002. Cell wall properties and their effects on the mechanical properties of fibers. J. Mater. Sci. 37 (1), 151
156. Boeckmann, B., Bairoch, A., Apweiler, R., Blatter, M.-C., Estreicher, A., Gasteiger, E., et al., 2003. The SWISS-PROT protein knowledge base and its supplement TrEMBL in 2003. Nucleic Acids Res. 31 (1), 365
370. Branden, C.I., 1999. Introduction to Protein Structure. Garland Science. Brodsky, B., Ramshaw, J.A., 1997. The collagen triple-helix structure. Matrix Biol. 15 (8), 545
554. Brown, R.M., 2004. Cellulose structure and biosynthesis: what is in store for the 21st century? J. Polym. Sci., A: Polym. Chem. 42 (3), 487
495. Cebe, P., Hu, X., Kaplan, D.L., Zhuravlev, E., Wurm, A., Arbeiter, D., et al., 2013. Beating the heat-fast scanning melts silk beta sheet crystals. Sci. Rep. 3, 1130. Cheng, S.Z., Cao, M., Wunderlich, B., 1986. Glass transition and melting behavior of poly (oxy-1,4-phenyleneoxy-1,4-phenylenecarbonyl-1,4-phenylene) (PEEK). Macromolecules 19 (7), 1868
1876.

128

Thermal Analysis of Textiles and Fibers

Choi, J., Kwak, S.-Y., 2007. Hyperbranched poly(ε-caprolactone) as a nonmigrating alternative plasticizer for phthalates in flexible PVC. Environ. Sci. Technol. 41 (10), 3763
3768. Crini, G., 2005. Recent developments in polysaccharide-based materials used as adsorbents in wastewater treatment. Prog. Polym. Sci. 30 (1), 38
70. Drummy, L.F., Farmer, B., Naik, R.R., 2007. Correlation of the β-sheet crystal size in silk fibers with the protein amino acid sequence. Soft Matter 3 (7), 877
882. Du, N., Liu, X.Y., Narayanan, J., Li, L., Lim, M.L.M., Li, D., 2006. Design of superior spider silk: from nanostructure to mechanical properties. Biophys. J. 91 (12), 4528
4535. Freedberg, I.M., Tomic-Canic, M., Komine, M., Blumenberg, M., 2001. Keratins and the keratinocyte activation cycle. J. Invest. Dermatol. 116 (5), 633
640. French, A.D., Johnson, G.P., 2009. Cellulose and the twofold screw axis: modeling and experimental arguments. Cellulose 16 (6), 959
973. Garel, A., Deleage, G., Prudhomme, J.-C., 1997. Structure and organization of the Bombyx mori sericin 1 gene and of the sericins 1 deduced from the sequence of the Ser 1B cDNA. Insect. Biochem. Mol. Biol. 27 (5), 469
477. Gelse, K., Po¨schl, E., Aigner, T., 2003. Collagens—structure, function, and biosynthesis. Adv. Drug Deliv. Rev. 55 (12), 1531
1546. Gomba´s, A., Szabo´-Re´ve´sz, P., Kata, M., Regdon, G., Er˝os, I., 2002. Quantitative determination of crystallinity of α-lactose monohydrate by DSC. J. Therm. Anal. Calorim. 68 (2), 503
510. Gosline, J.M., DeMont, M.E., Denny, M.W., 1986. The structure and properties of spider silk. Endeavour 10 (1), 37
43. Gosline, J., Guerette, P., Ortlepp, C., Savage, K., 1999. The mechanical design of spider silks: from fibroin sequence to mechanical function. J. Exp. Biol. 202 (23), 3295
3303. Grebowicz, J., Lau, S.F., Wunderlich, B., 1984. The thermal properties of polypropylene. J. Polym. Sci.: Polym. Symp. 71 (1), 19
37. Guille´n, F., Martı´nez, M.J., Gutie´rrez, A., Del Rio, J., 2005. Biodegradation of lignocellulosics: microbial, chemical, and enzymatic aspects of the fungal attack of lignin. Int. Microbiol. 8, 195
204. Gupta, R., Ramnani, P., 2006. Microbial keratinases and their prospective applications: an overview. Appl. Microbiol. Biotechnol. 70 (1), 21
33. Hatakeyama, T., Nakamura, K., Hatakeyama, H., 2000. Vaporization of bound water associated with cellulose fibres. Thermochim. Acta 352, 233
239. Hinman, M.B., Jones, J.A., Lewis, R.V., 2000. Synthetic spider silk: a modular fiber. Trends Biotechnol. 18 (9), 374
379. Holland, C., Vollrath, F., Ryan, A.J., Mykhaylyk, O.O., 2012. Silk and synthetic polymers: reconciling 100 degrees of separation. Adv. Mater. 24 (1), 105
109. Holland, P.M., Abramson, R.D., Watson, R., Gelfand, D.H., 1991. Detection of specific polymerase chain reaction product by utilizing the 50
30 exonuclease activity of Thermus aquaticus DNA polymerase. Proc. Natl. Acad. Sci. U.S.A. 88 (16), 7276
7280. Hotary, K., Allen, E., Punturieri, A., Yana, I., Weiss, S.J., 2000. Regulation of cell invasion and morphogenesis in a three-dimensional type I collagen matrix by membrane-type matrix metalloproteinases 1, 2, and 3. J. Cell Biol. 149 (6), 1309
1323. Hu, X., Kaplan, D., Cebe, P., 2006. Determining beta-sheet crystallinity in fibrous proteins by thermal analysis and infrared spectroscopy. Macromolecules 39 (18), 6161
6170. Hu, X., Kaplan, D., Cebe, P., 2007. Effect of water on the thermal properties of silk fibroin. Thermochim. Acta 461 (1), 137
144.

Thermal analysis of natural fibers

129

Hu, X., Kaplan, D., Cebe, P., 2008. Dynamic protein
water relationships during β-sheet formation. Macromolecules 41 (11), 3939
3948. Hu, X., Kaplan, D., Cebe, P., 2009a. Thermal analysis of protein
metallic ion systems. J. Therm. Anal. Calorim. 96 (3), 827
834. Hu, X., Lu, Q., Kaplan, D.L., Cebe, P., 2009b. Microphase separation controlled β-sheet crystallization kinetics in fibrous proteins. Macromolecules 42 (6), 2079
2087. Huang, J.-M., Yang, S.-J., 2005. Studying the miscibility and thermal behavior of polybenzoxazine/poly(ε-caprolactone) blends using DSC, DMA, and solid state 13 C NMR spectroscopy. Polymer (Guildf). 46 (19), 8068
8078. Huang, W., Krishnaji, S., Hu, X., Kaplan, D., Cebe, P., 2011. Heat capacity of spider silklike block copolymers. Macromolecules 44 (13), 5299
5309. Huo, P.P., Cebe, P., Capel, M., 1993. Dynamic mechanical relaxation and X-ray scattering study of poly(butylene terephthalate)/polyarylate blends. Macromolecules 26 (16), 4275
4282. Iervolino, E., Van Herwaarden, A., Van Herwaarden, F., Van de Kerkhof, E., Van Grinsven, P., Leenaers, A., et al., 2011. Temperature calibration and electrical characterization of the differential scanning calorimeter chip UFS1 for the Mettler-Toledo Flash DSC 1. Thermochim. Acta 522 (1), 53
59. Jiang, P., Liu, H., Wang, C., Wu, L., Huang, J., Guo, C., 2006. Tensile behavior and morphology of differently degummed silkworm (Bombyx mori) cocoon silk fibres. Mater. Lett. 60 (7), 919
925. Jin, H.-J., Kaplan, D.L., 2003. Mechanism of silk processing in insects and spiders. Nature 424 (6952), 1057
1061. Kadler, K.E., Baldock, C., Bella, J., Boot-Handford, R.P., 2007. Collagens at a glance. J. Cell Sci. 120 (12), 1955
1958. Kaplan, D., 1994. Silk Polymers: Materials Science and Biotechnology. Amer Chemical Society. Knight, D., Knight, M., Vollrath, F., 2000. Beta transition and stress-induced phase separation in the spinning of spider dragline silk. Int. J. Biol. Macromol. 27 (3), 205
210. Lau, S.F., Suzuki, H., Wunderlich, B., 1984. The thermodynamic properties of polytetrafluoroethylene. J. Polym. Sci.: Polym. Phys. Ed. 22 (3), 379
405. Lee, K., Ha, W., 1999. DSC studies on bound water in silk fibroin/S-carboxymethyl kerateine blend films. Polymer (Guildf). 40 (14), 4131
4134. Lourdin, D., Coignard, L., Bizot, H., Colonna, P., 1997. Influence of equilibrium relative humidity and plasticizer concentration on the water content and glass transition of starch materials. Polymer (Guildf). 38 (21), 5401
5406. Lu, S.X., Cebe, P., 1996. Effects of annealing on the disappearance and creation of constrained amorphous phase. Polymer (Guildf). 37 (21), 4857
4863. Ma, X., Yu, J., Kennedy, J.F., 2005. Studies on the properties of natural fibers-reinforced thermoplastic starch composites. Carbohydr. Polym. 62 (1), 19
24. Magin, T.M., Vijayaraj, P., Leube, R.E., 2007. Structural and regulatory functions of keratins. Exp. Cell Res. 313 (10), 2021
2032. Mathot, V., Pyda, M., Pijpers, T., Poel, G.V., Van de Kerkhof, E., Van Herwaarden, S., et al., 2011. The Flash DSC 1, a power compensation twin-type, chip-based fast scanning calorimeter (FSC): first findings on polymers. Thermochim. Acta 522 (1), 36
45. Mazzi, S., Zulker, E., Buchicchio, J., Anderson, B., Hu, X., 2014. Comparative thermal analysis of Eri, Mori, Muga, and Tussar silk cocoons and fibroin fibers. J. Therm. Anal. Calorim. 116 (3), 1337
1343.

130

Thermal Analysis of Textiles and Fibers

McGrath, K., Kaplan, D., 1997. Protein-Based Materials. Springer Science & Business Media. Menczel, J., Jaffe, M., 2007. How did we find the rigid amorphous phase? J. Therm. Anal. Calorim. 89 (2), 357
362. Menczel, J., Wunderlich, B., 1981. Heat capacity hysteresis of semicrystalline macromolecular glasses. J. Polym. Sci.: Polym. Lett. Ed. 19 (5), 261
264. Menczel, J., Wunderlich, B., 1986. Glass transition of semicrystalline macromolecules. In: Paper Presented at the Meeting of the Division of Polymer Chemistry, American Chemical Society. Menczel, J.D., Prime, R.B., 2014. Thermal Analysis of Polymers: Fundamentals and Applications. John Wiley & Sons. Miles, M.J., Morris, V.J., Orford, P.D., Ring, S.G., 1985. The roles of amylose and amylopectin in the gelation and retrogradation of starch. Carbohydr. Res. 135 (2), 271
281. Mizuno, M., Fujisawa, R., Kuboki, Y., 2000. Type I collagen-induced osteoblastic differentiation of bone-marrow cells mediated by collagen-α2β1 integrin interaction. J. Cell Physiol. 184 (2), 207
213. Moll, R., Divo, M., Langbein, L., 2008. The human keratins: biology and pathology. Histochem. Cell Biol. 129 (6), 705
733. Mora´n, J.I., Alvarez, V.A., Cyras, V.P., Va´zquez, A., 2008. Extraction of cellulose and preparation of nanocellulose from sisal fibers. Cellulose 15 (1), 149
159. Mullis, K., Faloona, F., Scharf, S., Saiki, R.K., Horn, G.T., Erlich, H. 1986. Specific enzymatic amplification of DNA in vitro: the polymerase chain reaction. In Cold Spring Harbor symposia on quantitative biology (Vol. 51, pp. 263
273). Cold Spring Harbor Laboratory Press. Mullis, K.B., Ferre´, F., Gibbs, R.A., 1994. The Polymerase Chain Reaction. Birkhauser Boston Inc. Murphy, A.R., Kaplan, D.L., 2009. Biomedical applications of chemically-modified silk fibroin. J. Mater. Chem. 19 (36), 6443
6450. Myllyharju, J., Kivirikko, K.I., 2004. Collagens, modifying enzymes and their mutations in humans, flies and worms. Trends Genet. 20 (1), 33
43. Omenetto, F.G., Kaplan, D.L., 2010. New opportunities for an ancient material. Science 329 (5991), 528
531. Park, S.J., Lee, K.Y., Ha, W.S., Park, S.Y., 1999. Structural changes and their effect on mechanical properties of silk fibroin/chitosan blends. J. Appl. Polym. Sci. 74 (11), 2571
2575. Pe´rez-Rigueiro, J., Viney, C., Llorca, J., Elices, M., 1998. Silkworm silk as an engineering material. J. Appl. Polym. Sci. 70 (12), 2439
2447. Piorkowska, E., Galeski, A., Haudin, J.-M., 2006. Critical assessment of overall crystallization kinetics theories and predictions. Progress Polym. Sci. 31 (6), 549
575. Pyda, M., Hu, X., Cebe, P., 2008. Heat capacity of silk fibroin based on the vibrational motion of poly(amino acid)s in the presence and absence of water. Macromolecules 41 (13), 4786
4793. Qin, G., Hu, X., Cebe, P., Kaplan, D.L., 2012. Mechanism of resilin elasticity. Nat. Commun. 3, 1003. Rajkhowa, R., Hu, X., Tsuzuki, T., Kaplan, D.L., Wang, X., 2012. Structure and biodegradation mechanism of milled Bombyx mori silk particles. Biomacromolecules 13 (8), 2503
2512. Ramiah, M., 1970. Thermogravimetric and differential thermal analysis of cellulose, hemicellulose, and lignin. J. Appl. Polym. Sci. 14 (5), 1323
1337.

Thermal analysis of natural fibers

131

Ricci, G., Patrizi, A., Bendandi, B., Menna, G., Varotti, E., Masi, M., 2004. Clinical effectiveness of a silk fabric in the treatment of atopic dermatitis. Br. J. Dermatol. 150 (1), 127
131. Rose, P., 1971. Protein-metal ion binding site: determination with proton magnetic resonance spectroscopy. Science 171 (3971), 573
574. Saiki, R.K., Scharf, S., Faloona, F., Mullis, K.B., Horn, G.T., Erlich, H.A., et al., 1985. Enzymatic amplification of beta-globin genomic sequences and restriction site analysis for diagnosis of sickle cell anemia. Science 230 (4732), 1350
1354. Samayam, I.P., Hanson, B.L., Langan, P., Schall, C.A., 2011. Ionic-liquid induced changes in cellulose structure associated with enhanced biomass hydrolysis. Biomacromolecules 12 (8), 3091
3098. Santin, M., Motta, A., Freddi, G., Cannas, M., 1999. In vitro evaluation of the inflammatory potential of the silk fibroin. J. Biomed. Mater. Res. 46 (3), 382
389. Schawe, J.E., 2014. Influence of processing conditions on polymer crystallization measured by fast scanning DSC. J. Therm. Anal. Calorim. 116 (3), 1165
1173. Scheller, J., Gu¨hrs, K.-H., Grosse, F., Conrad, U., 2001. Production of spider silk proteins in tobacco and potato. Nat. Biotechnol. 19 (6), 573
577. Schick, C., 2009. Differential scanning calorimetry (DSC) of semicrystalline polymers. Analyt. Bioanalyt. Chem. 395 (6), 1589
1611. Schultz, J.M., 2001. Polymer Crystallization: The Development of Crystalline Order in Thermoplastic Polymers. Amer Chemical Society. Schulz, E., Wunderlich, B., 1976. Macromolecular Physics, Vol. 2 Crystal Nucleation, Growth, Annealing, 50. Academic, New York. ˇ ´ k, J., Berggren, G., 1971. Study of the kinetics of the mechanism of solid-state reactions Sesta at increasing temperatures. Thermochim. Acta 3 (1), 1
12. Shao, Z., Vollrath, F., 2002. Materials: Surprising strength of silkworm silk. Nature 418 (6899), 741. Slade, L., Levine, H., 1995. Water and the glass transition—dependence of the glass transition on composition and chemical structure: special implications for flour functionality in cookie baking. J. Food Eng. 24 (4), 431
509. Suzuki, H., Grebowicz, J., Wunderlich, B., 1985. Heat capacity of semicrystalline, linear poly (oxymethylene) and poly (oxyethylene). Die Makromol. Chem. 186 (5), 1109
1119. Tajvidi, M., Falk, R.H., Hermanson, J.C., 2006. Effect of natural fibers on thermal and mechanical properties of natural fiber polypropylene composites studied by dynamic mechanical analysis. J. Appl. Polym. Sci. 101 (6), 4341
4349. Termonia, Y., 1994. Molecular modeling of spider silk elasticity. Macromolecules 27 (25), 7378
7381. Thygesen, A., Oddershede, J., Lilholt, H., Thomsen, A.B., Sta˚hl, K., 2005. On the determination of crystallinity and cellulose content in plant fibres. Cellulose 12 (6), 563
576. Tsuda, H., Kobayashi, M., Tanaka, T., Inoue, S., Magoshi, Y., & Magoshi, J. (2002). Different behavior of elution profile of silk fibroin on existence of calcium ion and potassium ion. In: Paper Presented at the Abstracts of Papers of the American Chemical Society. Turi, E., 2012. Thermal Characterization of Polymeric Materials. Elsevier. Uhlmann, D., 1972. A kinetic treatment of glass formation. J. Non-Crystal. Solids 7 (4), 337
348. Uitto, J., 2008. The role of elastin and collagen in cutaneous aging: intrinsic aging versus photoexposure. J. Drugs Dermatol.: JDD 7 (2 Suppl.), s12
s16.

132

Thermal Analysis of Textiles and Fibers

Van Herwaarden, S., Iervolino, E., Van Herwaarden, F., Wijffels, T., Leenaers, A., Mathot, V., 2011. Design, performance and analysis of thermal lag of the UFS1 twin-calorimeter chip for fast scanning calorimetry using the Mettler-Toledo Flash DSC 1. Thermochim. Acta 522 (1), 46
52. Varga, J., Menczel, J., Solti, A., 1984. Calorimetric study of the crystallization and melting of polypropylene. Polym. Sci. USSR 26 (12), 2763
2773. Vepari, C., Kaplan, D.L., 2007. Silk as a biomaterial. Progress Polym. Sci. 32 (8), 991
1007. Wang, X., Yucel, T., Lu, Q., Hu, X., Kaplan, D.L., 2010. Silk nanospheres and microspheres from silk/PVA blend films for drug delivery. Biomaterials. 31 (6), 1025
1035. Wang, Y., Kawano, Y., Aubuchon, S.R., Palmer, R.A., 2003. TGA and time-dependent FTIR study of dehydrating Nafion
Na membrane. Macromolecules 36 (4), 1138
1146. Wilkie, C.A., 1999. TGA/FTIR: an extremely useful technique for studying polymer degradation. Polym. Degrad. Stab. 66 (3), 301
306. Wortmann, F.J., Stapels, M., Elliott, R., Chandra, L., 2006. The effect of water on the glass transition of human hair. Biopolymers 81 (5), 371
375. Wu, Y.-J., Rheinwald, J.G., 1981. A new small (40 kd) keratin filament protein made by some cultured human squamous cell carcinomas. Cell 25 (3), 627
635. Wunderlich, B., 1976. Crystal Nucleation, Growth, Annealing. Academic Press. Xie, F., Yu, L., Chen, L., Li, L., 2008. A new study of starch gelatinization under shear stress using dynamic mechanical analysis. Carbohydr. Polym. 72 (2), 229
234. Xu, M., Lewis, R.V., 1990. Structure of a protein superfiber: spider dragline silk. Proc. Natl. Acad. Sci. U.S.A. 87 (18), 7120
7124. Xu, J.-T., Fairclough, J.P.A., Mai, S.-M., Ryan, A.J., Chaibundit, C., 2002. Isothermal crystallization kinetics and melting behavior of poly(oxyethylene)-b-poly(oxybutylene)/poly (oxybutylene) blends. Macromolecules 35 (18), 6937
6945. Yanagisawa, H., Davis, E.C., Starcher, B.C., Ouchi, T., Yanagisawa, M., Richardson, J.A., et al., 2002. Fibulin-5 is an elastin-binding protein essential for elastic fibre development in vivo. Nature 415 (6868), 168
171. Yokota, T., Oritani, K., Takahashi, I., Ishikawa, J., Matsuyama, A., Ouchi, N., et al., 2000. Adiponectin, a new member of the family of soluble defense collagens, negatively regulates the growth of myelomonocytic progenitors and the functions of macrophages. Blood 96 (5), 1723
1732. Young, P., Dollimore, D., Schall, C., 2000. Thermal analysis of solid-solid interactions in binary mixtures of alkylcyclohexanes using DSC. J. Therm. Anal. Calorim. 62 (1), 163
171. Zakzeski, J., Bruijnincx, P.C., Jongerius, A.L., Weckhuysen, B.M., 2010. The catalytic valorization of lignin for the production of renewable chemicals. Chem. Rev. 110 (6), 3552
3599. Zheng, Y., Bai, H., Huang, Z., Tian, X., Nie, F.-Q., Zhao, Y., et al., 2010. Directional water collection on wetted spider silk. Nature 463 (7281), 640
643. Zhou, L., Chen, X., Shao, Z., Huang, Y., Knight, D.P., 2005a. Effect of metallic ions on silk formation in the mulberry silkworm, Bombyx mori. J. Phys. Chem. B 109 (35), 16937
16945. Zhou, L., Terry, A., Huang, Y., Shao, Z., Chen, X., 2005b. Metal element contents in silk gland and silk fiber of Bombyx mori silkworm. Acta Chim. Sin.-Chin. Ed. 63 (15), 1379.

Polyester fibers

8

Michel Jaffe1, Anthony J. Easts2 and Xianhong Feng3 1 New Jersey Innovation Institute, University Heights, Newark, NJ, United States, 2 Consultant, Madison, NJ, United States, 3Beckton Dickinson and Company, Franklin Lakes, NJ, United States

Abstract Polyester fiber, specifically poly(ethylene terephthalate) (PET) fiber, is the largest volume synthetic fiber produced worldwide. The total volume produced in 2016 exceeded 50 million tons with a rate of growth far greater than any other fiber, natural or synthetic. Low cost, convenient processability, ease of blending with cotton and other natural fibers, convenient recyclability, and excellent and tailorable performance are the reasons for the dominating success of PET fiber. The excellent performance of polyester fiber over a wide range of end uses results from the ability to accurately control fiber morphology (distribution and connectivity of crystalline and noncrystalline load-bearing units) allowing the balance of thermal and dimensional stability, transport, and mechanical properties to be precisely controlled. All these parameters are conveniently and accurately monitored by thermal analysis techniques. It is the purpose of this chapter to provide the reader with an overview of the applications of thermal analysis toward polyester fiber characterization, including the impact of processing on performance and the utility of thermal analysis toward understanding the materials science of PET fibers.

8.1

Introduction

Polyester fiber, specifically poly(ethylene terephthalate) (PET) fiber, is the largest volume synthetic fiber produced worldwide. The total volume produced in 2016 exceeds 50 million tons with a rate of growth far greater than any other fiber, natural or synthetic. The distribution of synthetic fiber production by chemistry is shown in Fig. 8.1 (Industrievereingigung Chemiefase, 2018). If one assumes the total production is a single 5 dpf (B20 μm diameter) filament, the total length would be measured in light years (B1016 m) or the distance equivalent of leaving our solar system. While other polyesters are commercially produced in fiber form—polyethylene naphthalate (PEN), poly(butylene terephthalate) (PBT), poly(propylene terephthalate) (PPT), poly(lactic acid), and thermotropic polyester (liquid crystalline polymer—see chapter 19, Thermal analysis of liquid crystalline polymers)—these are of insignificant volume compared to PET; hence, Thermal Analysis of Textiles and Fibers. DOI: https://doi.org/10.1016/B978-0-08-100572-9.00008-2 © 2020 Elsevier Ltd. All rights reserved.

134

Thermal Analysis of Textiles and Fibers

Figure 8.1 Production of Polyester Fiber from 1975  2017 (in 1000 metric tons) (Industrievereingigung Chemiefase, 2018).

this chapter will focus primarily on PET. Recently, poly(ethylene furanoate)—PEF (the polyester of ethylene glycol and 2,5 furandicarboxylic acid), has been discussed as a commercial competitor to PET, but the emphasis is more on food packaging than fiber. The reasons for the dominating success of PET fiber are as follows: G

G

G

G

G

Low cost Convenient processability Ease of blending with cotton and other natural fibers Convenient recyclability Excellent and tailorable performance

The origin of low-cost polyester fiber lies in the efficient conversion of mixed xylenes to terephthalic acid; for details, see Modern Polyesters: chemistry and technology of polyesters and polyesters (Scheirs and Long, 2005). Cost and desirability is also positively impacted by the 280
C melting temperature of high-molecular weight PET, which allows the use of commercial heating fluids for processing, and the 75
C glass transition enables the morphology and molecular orientation introduced during processing to be stable at room temperature through washing temperatures. The excellent performance of polyester fiber over a wide range of end uses results from the ability to accurately control fiber morphology (distribution and connectivity of crystalline and noncrystalline load-bearing units), allowing the balance of thermal and dimensional stability, and transport and mechanical properties to be precisely controlled. All these parameters are conveniently and accurately monitored by thermal analysis techniques (Jaffe et al., 1981, 1997). By the end of the 20th century, PET fiber manufacture had shifted from the United States and Europe to Asia, with China and India dominating PET fiber manufacture, reducing the study of PET in the west but catalyzing research in the east. The quest for increased product sustainability has fostered the commercial development of PPT in the United States (see SCI, 1999; Scheirs and Long, 2005) and, more recently, PEF in the United States and Europe. Polyester fiber technology and performance have been reviewed in many publications (Eichhorn et al., 2009) and the reader is

Polyester fibers

135

directed to these publications for additional detail. It is the purpose of this chapter to provide the reader with an overview of the applications of thermal analysis toward polyester fiber characterization, including the impact of processing on performance and the utility of thermal analysis toward understanding the materials science of PET fibers.

8.2

Poly(ethylene terephthalate) history

The development of PET fiber, as with all synthetic thermoplastic fibers, originates with the research into aliphatic condensation polymers led by W. H. Carothers of DuPont in the 1930s (McIntyre, 2005). Much improved fiber performance was achieved in the early 1940s by the British team of J. Rex Whinfield and J. T. Dickson (Cheremisinoff, 1989), who focused their work on the aromaticaliphatic polyester of terephthalic acid and ethylene glycol. Commercialization of PET was rapid after World War II with the introduction of Terylene in Great Britain by ICI and the introduction of Dacron in the United States by DuPont, and PET fiber successfully entered the textile market as both filament yarn and staple and in the industrial market as a rubber reinforcement filament yarn, primarily for use in the sidewalls of radial passenger-car tires. Key performance advantages were wash-and-wear characteristics, blendability with cotton in textiles, and high modulus, coupled with excellent modulus retention, in industrial applications. Continued progress in the efficient, low-cost production of terephthalic acid ensured the dominance of PET as the fiber of choice in most fiber applications.

8.3

Poly(ethylene terephthalate) polymerization

PET is the condensation product of terephthalic acid and ethylene glycol. The key to successful PET polymerization is monomer purity and the absence of moisture in the reaction vessel. PET polymerization has often been reviewed (Burghardt and Vom, 1974) and the reader is referred to the many journal articles and patents dealing with all aspects of polyester fiber production (see, e.g., Nichols et al., 1996; Paszun and Spychaj, 1997). The first stage of PET polymerization is, in essence, the production of bis(hydroxyethyl)terephthalate (BHET). In the direct esterification of terephthalic acid, the reaction HOOC 2 C6 H4 2 COOH 1 2HOCH2 CH2 OH ! HOCH2 CH2 OCO 2 C6 H4 2 COOCH2 CH2 OH 1 H2 O

(8.1)

results in a mixture of low amounts of free BHET with a variety of PET oligomers. Water removal, usually under high vacuum conditions, is critical to the ultimate achievement of high molecular weights. The reaction catalysts for the

136

Thermal Analysis of Textiles and Fibers

ester interchange reaction have been the subject of intense research for many years and many catalyst compositions may be found in the patent literature (Wilfong, 1961; French Patent, 1959; Easley, 1959). The introduction of ester interchange catalysts requires the killing of these catalysts later in the polymerization sequence as they are equally effective as depolymerization catalysts. The next step in PET polymerization is melt condensation to high molecular weight. In this reaction step an ester interchange reaction occurs between two molecules of BHET to split off a molecule of glycol, building polymer molecular weight. The reaction must be catalyzed, and antimony trioxide, Sb2O3, is almost universally the moiety of choice. Typical melt-polymerization temperatures are at or above 285
C and viscosities are in the order of 3000 P, making uniform stirring and the imparting of uniform shear history, across the polymerization mixture, difficult to achieve. There is a prolific patent literature describing variations and improvements to PET polymerization. For reviews, see the work of Eichhorn et al. (2009) and East (2004). After achieving target molecular weight, the polymer can be pelletized for subsequent melt spinning (batch processing) or fed directly into a spinning machine and converted to fiber (continuous processing)—if the spun fiber is fed directly to a draw frame, the process is known as continuous polymerization (CP) spin-draw. The molecular weight of PET pellets can be further increased through solid-state polymerization. In this process, thoroughly dried PET chip is first crystallized at about 160
C (to prevent the as-polymerized chip from sticking together) and then heated just below the melting point under high vacuum and extreme dryness to advance the molecular weight upwards to values of inherent viscosity (IV) of 95 (textile grade chip has an IV of about 0.65) (Callender, 1985; Sorenson et al., 2001). The effects of the thermal-history process of PET chip and fiber have been extensively studied and are conveniently monitored by thermal analysis techniques. Jaffe et al. (1997) have reviewed the thermal behavior of PET and described the expected response of PET to process history in detail.

8.4

Characterization of poly(ethylene terephthalate) chip

PET chip or representative samples of CP spin-draw polymer are conveniently characterized by molecular weight, cleanliness, and thermal behavior. Molecular weight is characterized by the polymer IV, usually in halogenated solvents, and while the IV is related to molecular weight by the MarkHouwink equation, this is seldom done in practice. Polymer cleanliness is measured microscopically (optical techniques, polarized light microscopy) and is often expressed in subjective units such as the average number of black specks or the number of gels per gram of polymer. Cleanliness of the polymer is critical, a particle or gel of only a few microns in diameter can be responsible for a catastrophic spinline interruption. Thermal parameters are conveniently monitored by differential scanning calorimetry

Polyester fibers

137

(DSC), allowing a quick assessment of diethylen eglycol (DEG) content, crystallinity, etc. (Jaffe et al., 1997; Eichhorn et al., 2009).

8.5

Poly(ethylene terephthalate) fiber processing

The melt spinning of PET has been extensively treated in the patent literature, somewhat less so in the open literature (Davis and Hill, 1982), although the chapters by East (2004), Bessey and Jaffe (Ward et al., 2000), and Reese (2003) are good introductions to the process. Here, we concentrate on how changes in the key process variables of spinline stress and temperature profile affect the morphology developed in fiber spinning and drawing, and in turn, how the morphology affects the resulting performance of the yarn as manifested in thermal analysis. Process issues have been discussed in Chapter 2, Fiber processstructureproperty relationships, and that discussion is directly relevant to polyester-fiber production. The key process and performance parameters will be linked to the thermal analysis methods best utilized to elucidate the level and origin of the parameter of interest. The extraordinarily high volume of PET fiber produced in a given manufacturing plant, coupled with the huge length of fiber per unit mass, places increased emphasis on the statistics of any measurement made on fiber samples. In reality a coefficient of variation of 6 10% or less for key parameters (diameter, tensile properties, shrinkage) is required for the yarn to be commercially acceptable. For example, changes in noncrystalline molecular chain orientation (as manifested in fiber shrinkage at Tg), which is directly related to dye uptake and variations as low as 6 5%, may result in unacceptable lack of color uniformity in dyed fabrics. The frequency of variation in processing is also critical; high-frequency changes that may be averaged over a critical yarn length, are, in general, more acceptable than smaller variation along or between specific filaments, which occurs at the lower frequency. The materials science of synthetic-fiber spinning discussed in Chapter 2, Fiber processstructureproperty relationships, defines how morphology develops during spinning and drawing and is directly applicable to PET fiber. The remainder of this chapter will focus on the specifics of the thermal analysis of PET. Specific process details tend to be less discussed in the open literature and the reader is thus directed to the patents of Celanese, DuPont, Fiber Industries, and Allied Chemical Corporation (none of these companies currently exist) and the more recent patents of Reliance, in India, and, for example, the Jiangsu Sanfangxiang Group, in China. Much of PET yarn production is converted to “staple” fiber and the demands of staple fiber are different to those of filament yarns. Staple fiber is a continuous filament cut into short lengths of B30100 1 mm. Staple fibers are discontinuous and are crimped and chopped to the desired staple-fiber length to effectively blend with cotton, wool, or other natural fibers. The raw polyester fibers are melt-spun through many hundred spinneret holes and collected in large drums or cans; fiber

138

Thermal Analysis of Textiles and Fibers

bundles from many cans are combined into thick bundles of fibers called a “tow,” which is often of several million decitex. These thick bundles of fibers are then drawn on a massively constructed draw frame, heat-set in a steam-heated hot box, and then usually crimped using a “stuffer box” method. The bulked tow is finally cut to the desired staple length and compressed into bales. After spinning and drawing are completed, polyester yarns become components in a variety of end-use products with textiles (fabrics) and rubber reinforcements (tire cord) comprising the highest volume. Every end use involves subjecting the fiber to additional temperature and stress history, impacting the observed thermal response of the fibers being tested. This history is manifested in DSC, thermo mechanical analysis (TMA), and differential thermo mechanical analysis (DTMA) and examples are given in the discussions below.

8.6

Physical properties of poly(ethylene terephthalate)

PET is a semicrystalline polymer and its physical parameters have been repeatedly determined over many years. Table 8.1 is a summary of typically accepted values (Kitano et al., 1995).

8.7

Other polyesters

Other aromatic, aliphatic polyesters of commercial import include PEN, PPT, and PBT. These are produced in far smaller volumes than PET and are often focused toward specific markets: PEN for cordage, PPT for carpets, and PBT for stretched fabrics. PEF fiber (patents go back over 60 years) is similar in performance to PET but is not yet commercially produced. Initial applications for PEF are focused on food packaging, but if commercial introduction is successful, the conversion of PEF to fiber products is likely.

Table 8.1 Physical Parameters of PET. Crystal habit Cell parameters Cell density Tm (DSC) ΔHf Tg (solid chip) Tg (drawn fiber) Specific gravity

Triclinic: one polymer chain per unit cell a 5 0.444 nm; b 5 0.591 nm; c 5 1.067 nm, α 5 100 degrees; β 5 117 degrees; γ 5 112 degrees 1.52 g/cm3 260
C265
C 140 J/g; 33.5 cal/g 79
C (DSC) 120
C (dynamic loss) 1.33 (amorphous, undrawn), 1.39 (crystalline drawn fiber)

Polyester fibers

8.8

139

Thermal analysis (TA) of polyester fibers

The Thermal analysis (TA) literature associated with PET fiber is extensive and has been thoroughly reviewed by Jaffe (Jaffe et al., 1981, 1997), and others (Eichhorn et al., 2009) from the inception of modern thermal analysis in the 1960s to about 2010. The physical parameters of PET (TgB79
C, TmB175
C) and their molecular relaxation times and crystallization rates enable convenient control of fiber morphology and allow the production of fibers for a broad range of applications. Modern PET is usually spun at very high speeds of $ 5000 m/min; so, examples of thermal analysis results that show the impact of spinning speed (spinline stress) will be emphasized. Fig. 8.2 shows typical PET DSC behavior with Tg, crystallization of heating (cold crystallization), melting, crystallization upon cooling, and, if heated above B325
C, decomposition. All phenomena associated with the physical changes, which are known to occur in semicrystalline polymers, are noted in DSC studies of PET. Changes in process history manifest in the location, size, and shape or the resulting DSC traces. For example, changes in the melting of as-spun PET, as a function of spinning stress (or spinning speed, molecular weight (MW), etc.), is shown in the classic work of Heuvel and Huisman (1978) in Fig. 8.3. Fig. 8.3 clearly shows that as the spinning speed (spinning stress), applied to the spinning PET fiber, is increased, the temperature of cold crystallization goes down toward Tg and the spinning PET fiber decreases in size, while the crystallinity of the spun yarn increases. Fig. 8.4 shows the premelting observed in PET fibers as a function of yarn thermal-history (annealing temperature). The premelt endotherm represents the

Figure 8.2 DSC of as-spun PET fiber (A) is the first heating of a typical as-spun fiber, (B) is the first cooling after melting, (C) is the second heating (Ms. Cindy Lee, Ms. Berta Marani, Dr. Michael Jaffe, New Jersey Innovation Institute, unpublished results). PET, Poly (ethylene terephthalate).

Figure 8.3 DSC of PET spun yarn as a function of spinning speed. PET, Poly(ethylene terephthalate). Source: From Heuval and Huisman (1978), J. Appl. Polym. Sci. 22(8), 22292243, reprinted with permission from John Wiley & Sons, Ltd.

Figure 8.4 “Premelt” endotherm of commercial PET fiber as a function of thermal history. PET, Poly(ethylene terephthalate). Source: From Berndt and Bossman (1976), Polymer 17(3), 241245, reprinted with permission of Butterworth Heinemann, Ltd.

Polyester fibers

141

melting of crystals formed during a heat-treatment step of the fiber or fabric; these crystals are stable up to the temperature of formation and provide useful diagnosticprocess information. Such studies, pioneered by the late Edith Turi on nylon fibers, give useful insight into the time and temperature history of polyester fibers. Monitoring the dimensional stability (TMA) and mechanical properties of polyester fiber yields complementary information of DSC and provides insight into the stress history of the fibers under investigation while simultaneously providing data of end-use significance (shrinkage at a given temperature). Fig. 8.5 shows a diagrammatic representation of the dimensional change, occurring as an oriented and crystallized PET fiber, which is allowed to shrink in a close-to-zero tension fiber experiment. Process history affecting behavior at Tg, Tm, and processing temperatures manifests in the resulting plots of dimensional change versus temperature. TMA can be utilized to measure the coefficient of linear expansion (CTE) in the direction of the long axis of the fiber. It has been established that as the molecular orientation of a fiber increases, the CTE decreases from the isotropic value of CVE/ 3 to the change in the crystal unit cellchain dimension change in the fiber direction. As illustrated in Fig. 8.6, when measured in a region of temperature below Tg, a useful way is provided to monitor average molecular orientation in the fiber, which is especially useful if the fiber cross-section is not circular. Examination of the dimensional stability changes in as-spun PET yarns clearly shows the structural changes occurring as the line speed (spinning stress) is increased. Referring again to the classic work of Heuval and Huisman (1978), it is shown in Fig. 8.7, showing first the increase of molecular orientation with increasing spinning speed, accompanied by an increase in fiber initial modulus and, as higher spinning speed, the introduction of crystallinity in the spinline, shown by the shifting of shrinkage from Tg to Tm. Comparing Fig. 8.7 to Fig. 8.3 illustrates the

Figure 8.5 Diagrammatic representation of the dimension changes in processed PET fiber as a function of process history (Jaffe, Celanese Research Company, unpublished results). PET, Poly(ethylene terephthalate).

142

Thermal Analysis of Textiles and Fibers

Figure 8.6 Coefficient of linear thermal expansion of a PET fiber as a function of average molecular orientation (Jaffe, Celanese Research Company, unpublished results). PET, Poly (ethylene terephthalate).

Figure 8.7 Shrinkage of as-spun PET yarn as a function of spinning speed (Heuval and Huisman, 1978). PET, Poly(ethylene terephthalate). Source: Reprinted with permission of John Wiley & Sons.

Polyester fibers

143

detailed process understanding that can be obtained when the thermal response of variously processed fibers is examined by a variety of complementary techniques. Dynamic mechanical analysis of polyester fibers yields useful end-use characterization of polyester fibers by providing information to fiber modulus as a function of temperature as well as a measure of Tg and a measure of the work loss exhibited by the fiber during deformation. Fig. 8.8 represents the typical DMA data on PET

Figure 8.8 DMA of a drawn PET fiber as a function of the draw ratio (Miller and Murayama, 1984). PET, Poly(ethylene terephthalate). Source: Reprinted with permission of John Wiley & Sons, Ltd.

144

Thermal Analysis of Textiles and Fibers

fibers as a function of the applied draw ratio, showing the increase in initial fiber modulus and the change in shape of tan δ. The examples shown before illustrate how the thermal analysis of PET fiber reflects the details of the fiber morphology and can be related to process history and end-use performance. An excellent illustration on how this data may be used is illustrated in the early work of the Valk et al. (1980) shown in Fig. 8.9, which shows the relationship of fabric dye uptake to the annealing stress and temperature applied to the fabric during heat setting (annealing under tension). One could substitute fabric shrinkage at a given temperature or the amorphous molecular orientation function for the dye uptake axis with similar results. If the fabric was made of a polyester fiber with a different process history, the details of the plot would change to reflect these differences. The examples shown illustrate the power of thermal analysis in the characterization of PET fibers. Inherent in these examples are the several ways in which the similar PET fiber thermal data can be utilized by the fiber scientist: 1. Finger printing: Does the fiber, produced today, match the thermal characteristics of fiber produced yesterday and does it also meet the product specifications? 2. Quality control: Do the end-use properties, such as shrinkage at a given temperature (TMA), coefficient of thermal expansion along the fiber axis (TMA), initial modulus or hysteresis as a function of stretch/relax cycles (DMA), or overall structural stability as measured by crystallinity (DSC) meet product specifications?

Figure 8.9 Relationship of fabric dye uptake to fabric heat-setting process history. Source: From Valk et al. (1980), Textile Research Journal, 50(1), 4654, reprinted with permission of Deutcher Verlag Gmbh.

Polyester fibers

145

3. Process development/diagnostic: A highly efficient tool to assist new product development and/or a diagnostic for what is causing a given process to yield off-specification materials.

Similar techniques and experimental approaches can be applied to other aromaticaliphatic polyesters of fiber importance including PEN, PPT, PBT, and perhaps in the future, PEF.

8.9

Polyester fiber thermal analysis in the 21st century

As polyester fiber products became mature, studies of the thermal behavior of PET and other polyester fibers became less frequent and tended to focus on specialized rather than general aspects of polyester fibers. In a major study of PET industrial yarns, Liu et al. (2016) compared the morphology of HMLS (high modulus low shrinkage) yarn to that of HMLE (high modulus low elongation). HMLS is produced from a highly oriented precursor yarn, while HMLE begins with low orientation spun yarn; both are drawn to maximum draw ratio and heat-set (details of processing varies with individual producers). Liu (2016) employed wide angle x-ray scattering (WAXS), small angle x-ray scattering (SAXS), and DSC for determining fiber morphology. They conclude that HMLS has higher crystallinity and more perfect crystals than HMLE. Comparison of the DSC traces of HMLS and HMLE, shown in Fig. 8.10, supports these conclusions, although the detailed X-ray study performed provided more detailed morphological

Figure 8.10 Comparison of the melting behavior of HMLS and HMLE PET industrial yarns (Liu et al., 2016). HMLE, High modulus low elongation; HMLS, high modulus low shrinkage; PET, poly(ethylene terephthalate).

146

Thermal Analysis of Textiles and Fibers

information and supports the importance of the use of a variety of techniques when studying the details of fiber structure. Yoon (Yoon et al., 2017) showed how a modified PET drawing process increased the crystallinity of PET industrial yarns and utilized DSC to measure crystal-related changes. Potter (2012) published a review on the use of DSC and TGA for the identification of textile fibers, including PET and other polyesters. Gashti (Gashti and Navid, 2012) used DSC, TGA, and dynamic mechanical thermal analysis (DMTA) to study the effect of silicone emulsion softeners on the thermal and flammability properties of the resulting textile fibers. Their findings show that the application of the silicone coating leads to an increased rate of thermal degradation and increased flammability of the resulting textile. Telli and Kale (2011) examined the effect of ZnO nanoparticles, added to impart antibacterial properties to the PET fiber, on the thermal properties of the resulting fiber. Fig. 8.11 shows the changes in crystallization behavior as monitored by DSC, proving that the ZnO is acting as a nucleating agent for the PET, as evidenced by the increase of the crystallization temperature as a function of ZnO content. In a DSC and TGA study of polyester containing a Zn ion and phosphinic acid compound flame retardant, Liu et al. concluded that the flame retardant enhanced

Figure 8.11 Changes in PET fiber as a function of ZnO concentration, MB signifies the PET master batch used to spin fibers (Telli and Kale, 2011). PET, Poly(ethylene terephthalate).

Polyester fibers

147

the thermal stability of the resulting resin (Liu et al., 2014). Chen et al. (2014) studied the thermal conversion mechanism of polyester fiber (Dacron) and a variety of other materials by TGA. A study of PET fiber cross-linking by trimethlylolpropane and electron beam radiations was studied by DSC and TGA and the impact on structure and properties was monitored (Zhu et al., 2017). In contrast to the fundamental studies of the materials science of PET fiber in the last century, work in the 21st century highlights the utility of thermal analysis in the study of specific PET fiber products. Although the focus of the work has changed, the use of thermal analysis techniques to elucidate the chemical and morphological impact on polyester fiber performance continues to be a basic tool in the study of presently ubiquitous polyester fiber products.

8.10

Other polyester fibers, polypropylene terephthalate, poly(butylene terephthalate), polyethylene naphthalate

Most of the work done on the thermal analysis of polyesters, other than PET, is focused more on plastic than on fiber applications. In an article published in the online New Fibres magazine in 2000, Brown et al. (2000) show TMA data for PTT spunbond fibers. Thermal analysis of PEN fibers was treated in detail by Menczel in 1996 (Jaffe et al., 1997). In general, any data relevant to molecularly oriented structure of PPT, PBT, or PEN will relate to the expected thermal behavior of these polymers in fiber form.

8.11

Conclusion

Polyester fiber thermal analysis has been extensively studied over the past 50 years, resulting in a rich literature that utilizes thermal analysis techniques to develop processstructureproperty relationships as well as to measure properties directly related to polyester fiber and fiber-based product performance. The power of thermal analysis to provide a basic understanding of the mechanisms and processes, which allow for the huge volume production of commercially important fiber products, cannot be overstated.

References Berndt, H.-J., Bossman, A., 1976. Polymer 17 (3), 241245. Brown, H., Casey, P., Donahue, J., 2000. NF New Fibres Magazine. On-line.

148

Thermal Analysis of Textiles and Fibers

Burghardt, W., Vom, O.H., 1974. Process for the Manufacture of Fibers from High Molecular Weight Linear Polyethylene Terephthalate, US 3803284 A. Callender, D.G., 1985. Polym. Sci. Eng. 25, 453457. Chang, Y., et al., 2004. Macromol. Mater. Eng. 289 (8), 703707. Chen, et al., 2014. Fuel 137, 7784. Cheremisinoff, N.P., 1989. Handbook of Polymer Science and Technology, Volume 1: Synthesis and Properties. Marcel Dekker. Davis, G.W., Hill, E.E., 1982. Kirk-Othmer Encyclopedia of Chemical Technology, Third ed. Wiley, p. 535. Easley, W.K., (1959). Canadian Patent 573,301 to Chemstrand Corp. East, A.J., 2004. Kirk-Othmer Encyclopedia of Chemical Technology, vol. 5. Wiley. Eichhorn, S., Hearle, W.S., Jaffe, M., Kikutani, T., 2009. Handbook of Textile Fiber Structure, vol. 1. Woodhead Publishing, Oxford, Cambridge, New Delhi. French Patent, 1959. French Patent 1,169,659 to I.C.I. Gashti, M.P., Navid, M.Y., 2012. J. Appl. Polym. Sci. 125 (2), 14301438. Heuvel, H.M., Husiman, R., 1978. J. Appl. Polym. Sci. 22 (8), 22292243. Jaffe, M., Menczel, J.D., Bessey, W.E., 1981. Fibers. In: Turi, E. (Ed.), Thermal Characterization of Polymer Materials. Academic Press, New York. Jaffe, M., Menczel, J.D., Bessey, W.E., 1997. Fibers. In: Turi, E. (Ed.), Thermal Characterization of Polymer Materials, second ed. Academic Press, San Diego, CA. Kitano, et al., 1995. Polymer 36, 10. Liu, Y., et al., 2014. Fangzhi Xuebao 35, 1216. Liu, Y., et al., 2016. Polymer 105, 157166. McIntyre, J.E., 2005. Synthetic Fibers: Nylon, Polyester, Acrylic, Polyolefin. Woodhead Publishing. Miller, R.W., Murayama, T., 1984. J/ Appl. Polym. Sci. 29 (3), 933939. Nichols, C.S., Moore, T.C., Edwards, W.L., 1996. Method of Post-Polymerization Stabilization of High Activity Catalysts in Continuous Polyethylene Terephthalate Production, US Patent 5898058 A. Paszun, D., Spychaj, T., 1997. Ind. Eng. Chem. Res. 46 (4), 13731383. Potter, C., 2012. AATCC Rev. 12 (5), 3945. SCI, 1999. Process Economics Program Report, 227, SCI International Reports. Reese, G., 2003. Encyclopedia of Polymer Science and Technology, third ed. Wiley, pp. 652678. Scheirs, J., Long, T., 2005. Modern Polyesters: Chemistry and Technology of Polyesters and Copolyesters. Wiley. Sorenson, F., et al., 2001. Preparative Methods of Polymer Chemistry. Wiley. Statista, 2019. Industrieverereingung Chemiefase, August, 2018. Telli, M., Kale, R., 2011. Adv. Appl. Polym. Res. 2 (4), 491502. Valk, G., Jellinek, G., Schroder, U., 1980. Tex. Res. J 50 (1), 4654. Ward, I., et al., 2000. Solid Phase Processing of Polymers. Hansa Verlag. Wilfong, R.E., 1961. J. Polym. Sci. 54, 388. Yildirim, K., Ulcay, Y., 2014. e-Polymer 14 (2), 12. Yoon, J., et al., 2017. Polym. Sci. Eng. 57 (2), 1290. Zhu, S., Shi, M., Tian, M., Qu, L., Chen, G., 2018. Effects of irradiation on polyethyleneterephthalate (PET) fibers impregnated with sensitizer. J. Text. Inst. 109 (3), 294299.

Polyester fibers

149

Further reading Lewin, M., 2006. Handbook of Fiber chemistry, third ed. CRC Press. Menczel, J.D., Prime, B., 2009. Thermal Analysis of Polymers: Fundamentals and Applications. Wiley. Saw, C.K., Menczel, J.D., Choe, E. W., Hughes, O. R. (1997), SPE/ANTEC Proceedings, 610. Towe, T., 2009. Interior Textiles Design and Developments. Woodhead Publishing.

Thermal properties of aliphatic polyesters

9

Mazeyar Parvinzadeh Gashti1,2, Marlon Bustos1, Hicham Alayan1, Roya Jamarani1 and Milan Maric1 1 Department of Chemical Engineering, McGill University, Montre´al, QC, Canada, 2 Research and Development Laboratory, PRE Labs Inc., Kelowna, BC, Canada

Abstract Thermal properties of polyesters are one of the most important parameters in their manufacturing and processing in various industries. Studies have shown that the thermal properties of these polymers are very dependent to their chemical structure in different forms including aromatic, semi-aromatic and aliphatic chains. In addition, the properties of polyesters can be changed in homopolymeric and copolymeric forms. Different additives and fillers have been widely considered to reinforce polyester and increase the thermal stability of final composites. In this review, we presented research studies devoted to thermal properties of polyglycolides, polylactic acid/poly(lactide)s, polycaprolactones and polyhydroxyalkanoates.

9.1

An introduction to aliphatic polyesters

Polyesters are the most widely used polymers in textiles, resins, sheets, packaging, and medical materials (Koronis et al., 2013; Tan et al., 2013; Zia et al., 2016; Parvinzadeh Gashti and Moradian, 2012; Parvinzadeh and Hajiraissi, 2008; Ebrahimi et al., 2011). One representative polyester is poly(ethylene terephthalate) (PET), which is very common in the production of textile fibers (Parvinzadeh Gashti et al., 2011, 2013a, 2014, 2015; Parvinzadeh and Hajiraissi, 2007). Variants of polyesters are aromatic, semi-aromatic, and aliphatic (Braun and Levin, 1986). Common to all polyesters is the presence of ester linkages in their backbones which are derived most often from esterification of poly-functional alcohols and acids (Paul et al., 2015). Vectran is a commercial high-performance aromatic polyester produced by Celanese Acetate LLC with excellent thermal stability, mechanical properties, and chemical stability (Menczel et al., 1997). Examples of commercial semi-aromatic polyesters are PET, poly(butylene terephthalate), poly(trimethylene terephthalate), and poly(ethylene naphthalate). It has been demonstrated that the aromatic sections of polyesters generally contribute to improved thermal stability and mechanical properties (Nejad et al., 2012; Parvinzadeh Gashti et al., 2013b; Hajiraissi and Parvinzadeh, 2011). Thermal Analysis of Textiles and Fibers. DOI: https://doi.org/10.1016/B978-0-08-100572-9.00009-4 © 2020 Elsevier Ltd. All rights reserved.

152

Thermal Analysis of Textiles and Fibers

Aliphatic polyesters are some of the most well-known biodegradable polymers in the polyester family which are used in tissue engineering, medical science, and drug delivery (Vert, 2005). This is due to their unique degradation in tissues and biodegradability in the environment. The hydrolysis rate and functionality depend on the temperature, macromolecular structure, synthesis method, and density of ester domains (Vert, 2005; Tokiwa and Calabia, 2004). The biodegradability of these polymers is only achievable if all its organic components undergo a complete degradation by naturally occurring microorganisms, such as bacteria, fungi, and algae (Vert, 2005; Tokiwa and Calabia, 2004). Some of the most common biodegradable polyesters are polyglycolide (PGA), polylactide (PLA), polycaprolactone (PCL), and polyhydroxyalkanoates (PHAs). These aliphatic polyesters represent excellent candidates for biomedical materials, packages, and sheets in food and textile industries (Vert, 2005; Tokiwa and Calabia, 2004; Seyednejad et al., 2011). Aliphatic polyesters are mainly synthesized by polycondensation, ring-opening polymerization, and enzymatic polymerization (Seyednejad et al., 2011; Jerome and Lacomte, 2008). Briefly, polycondensation of aliphatic polyesters refers to stepwise condensation, typically of diols and diacids. Another effective polymerization method is based on enzymatic-activated monomers with the aid of lipase (Seyednejad et al., 2011; Jerome and Lacomte, 2008). Each of these methods has its own advantages and disadvantages which are not discussed in this chapter. As previously described, a significant effort has been devoted to the development of aliphatic polyesters in alternative industries, other than tissue engineering and medical devices. Attempting to fulfill requirements in these other applications necessitates a structural measurement, such as evaluation of thermal properties. A researcher needs to measure the thermal properties of polymers in order to understand the formulation, process history, and utility for target applications (Ozawa, 2000). Several parameters can be evaluated by thermal analytical tools including heat capacity, thermal expansion, thermal conductivity, and thermal shock resistance. In other words, melting and crystallization characteristics, thermal history, nucleation, crystallinity mechanisms, and degradation temperature of polymers can be studied through thermal analysis techniques. Thermal properties of polymers are often measured using differential scanning calorimetry (DSC), thermogravimetric analysis (TGA), and dynamic mechanical analysis (DMA) (Forrest et al., 1997; Dalkoni-Veress et al., 2001; Rahimi et al., 2011; Parvinzadeh Gashti et al., 2012, 2013c). Typically, DSC curves are obtained by heating polymer samples as a function of time or temperature at a constant rate of heating. From the plots, the traces are shifted due to the change in heat capacity of samples. As a result of the increased molecular motions in the polymer, the enthalpy changes can be calculated by endothermic or exothermic shifts, and several parameters can be detected, such as solid
solid transitions, glass transition, crystallization temperature, and melting temperature (Coleman and Craig, 1996). The thermal stability of polymers is often assessed by TGA and is usually the first step before performing DSC. This method generally refers to the energy needed to decompose the polymer as a function of increasing temperature at a

Thermal properties of aliphatic polyesters

153

constant heating rate. It has been shown that TGA curves of polymers basically consist of three regions: initial degradation, main degradation, and char decomposition. The first region of degradation occurs at relatively low temperatures during the heating of polymers, at which degradation occurs as a result of removing water molecules or desorption of the physically adsorbed solvents in polymer. The second region of degradation is generally the most significant, with the highest percentage of polymer mass loss, which occurs due to cleavage of the main polymer chains and crystalline structures. The production of char and final volatile products finally occurs in the third region, which involves dehydration and charring reactions and results in the release of water and carbon dioxide (Parvinzadeh Gashti and Parvinzadeh Gashti, 2012; Parvinzadeh Gashti et al., 2016). In DMA analysis, the elastic modulus of a polymer as well as mechanical properties is determined as a function of applied frequency and temperature. This method is very helpful to understand the viscoelastic and compositional properties of polymers. According to earlier studies, thermal properties of aliphatic polyesters generally depend on the synthesis method, the structure of main chain and side groups, degree of branching, degree of crystallinity, and the type and amount of additives used in blending (Savelyeva et al., 2015; Erythropel et al., 2016). The aim of this chapter is to provide an overview of the thermal properties of aliphatic polyesters and discuss the most recent techniques used to improve the thermal properties of these polymers, such as copolymerization, blending, and composite formation.

9.2

Thermal properties of polyglycolides

9.2.1 Homopolymers PGA is one of the most used thermoplastic biopolymers in medical and pharmaceutical industries with the simplest aliphatic structure of the polyester group (Fig. 9.1) (Ashammakhi, 1996). PGA can be produced through polycondensation, high temperature cationic ring-opening polymerization, and solid-state polycondensation of halogen acetates (Doppalapudi et al., 2014). Thermal measurement of PGA is important in terms of biodegradability, enzymatic, and hydrolytic degradation. Generally, a relatively high-molecular mass biopolymer should be synthesized to guarantee a sufficient decomposition and degradation time in tissue engineering and medical applications. This has a direct relationship to thermal degradation and stability of synthesized PGA (Ashammakhi, 1996; Doppalapudi et al., 2014).

Figure 9.1 The chemical structure of PGA biopolymer.

154

Thermal Analysis of Textiles and Fibers

Great effort has been devoted in the last two decades to study the thermal properties and biodegradability of PGA homopolymer. Mathias et al. introduced a onestep method to prepare PGA homopolymer by reacting bromoacetic acid and triethylamine in anhydrous diethyl ether at room temperature. They found the melting point (Tm) of the final product to be 196 C. Their method of synthesis was useful to cross link vinyl monomers for biodegradable future applications. Considering the fact that a highly porous PGA is needed for bone tissue regeneration (Ashammakhi, 1996), solid-state polymerization was used by Epple and Herzberg (1998) who produced PGAs with different MOOC
CH2
X precursors (M 5 metal, X 5 halides). Epple and Herzberg (1998) claimed that the melting enthalpies of PGA change from 3.2 to 7.1 kJ/mol, depending on the precursor and production conditions. However, they did not show their observations in detail. DSC results revealed that the final PGA had a crystallinity in the range of 30%
60%, with the appearance of an exothermic peak in the curve corresponding to residual byproducts. In another study, Epple and Herzberg (1998) achieved a relatively pure PGA with high-molecular mass due to the absence of diglycolide in the final material (Schwarz and Epple, 1999). This result was because they did not detect any glass transition (Tg), exothermic recrystallization (Tc), and endothermic melting peaks related to diglycolide side products (Fig. 9.2A). Based on TGA results for solidstate
produced PGA, the main thermal degradation of high-molecular mass and mass loss of polymer started at 220 C with two derivative thermogravimetric peaks at 317 C and 370 C, while small-chain oligomers degraded below 220 C (Schwarz and Epple, 1999). Their synthesis method was also utilized in another study to produce microporous PGAs containing the drug Goserelin with controlled release properties (Schwarz and Epple, 1999). On the other hand, Roushandeh and Nabi Sarbolouki (2001) improved the thermal stability of PGA at which their samples started to degrade at 272 C. Homopolymers were obtained by ring-opening polymerization of glycoside (GL) in the presence of a co-initiator and catalyst. Their synthesis method resulted in the production of high-molecular mass PGA in the presence of tin catalysts. This was due to the fact that they were able to precisely control the equilibrium nature of the reaction. Furthermore, they had no difficulties in removing water, which is commonly generated during the production of low-molecular mass PGAs, from the synthesized polymer (Roushandeh and Nabi Sarbolouki, 2001). An interesting study was conducted by You et al. (2005) to evaluate the impact of exposing electrospun PGA tissues to in-vitro biodegradation solutions on thermal parameters at different treatment conditions. As expected, the Tm peak was 222 C for the untreated sample while it shifted to 220 C, 207 C, and 164 C for 1, 8, and 12 days treatment, respectively, confirming that degradation started in the amorphous regions before expanding to the crystalline regions (Fig. 9.2B). Some researchers have been trying to develop PGA homopolymers with improved thermal properties by establishing new synthesis conditions. In this regard, the research done by Jing et al. (2007), Schmidt et al. (2014), and Baez and Marcos-Fernandez (2015) are worth mentioning. Jing et al. (2007) prepared cyclohexyl-substituted PGAs in solution form and in bulk form, aiming to increase

Thermal properties of aliphatic polyesters

155

Figure 9.2 (A) DSC traces of diglycolide (top) and solid-state chemically prepared polyglycolide (bottom). (Mettler TA 4000 instrument in sealed aluminum crucibles. The heating rate was 5K/min and the average sample mass was 5
8 mg. Figure 6 in original paper.) (B) DSC thermograms for electrospun PGA during in vitro degradation [DSC measurements were conducted with a Perkin
Elmer (MA) DSC7 instrument under a nitrogen atmosphere. About 5 mg samples was sealed in an aluminum pan for the measurements. The samples were heated from 10 C to 250 C at a rate of 20 C/min, held at 250 C for 1 min, and then quenched to 10 C. The samples were reheated to 250 C. Figure 7 (a) in original paper]. (C) DSC thermograms of PGA samples with different Mn values. [The thermal properties of the polymers were analyzed with a DSC 1/500658/200W STARe System from Mettler Toledo equipped with an FRS5-sensor and nitrogen cooling. For each sample, two heating and one cooling cycles were measured. Only the last heating cycle was analyzed to determine the thermal properties of the material. All samples were analyzed using a constant heating and cooling rate of 10K/min in a temperature range from 260 C to 280 C. Figure 7(a) in original paper.] (D) TGA curves of PGA samples with different Mn value [the thermal degradation of the polymers was measured with a Netzsch TGA 204 Phoenix (Selb, Germany) instrument at a constant heating rate of 10K/min in a temperature range from 25 C to 400 C under argon atmosphere. Figure 7(b) in original paper]. Source: (A) From Schwarz, K., Epple, M., 1999. Macromol. Chem. Phys. 200, 2221. Reprinted with permission from John Wiley and Sons; (B) From You, Y., Min, B.M., Lee, S. J., Lee, T.S., Park, W.H., 2005. J. Appl. Polym. Sci. 95, 193. Reprinted with permission from John Wiley and Sons; (C and D) From Schmidt, C., Behl, M., Lendlein, A., Beuermann, S., 2014. RSC Adv. 4, 35099. Reprinted with permission from the Royal Society of Chemistry.

156

Thermal Analysis of Textiles and Fibers

the Tg. They interestingly found that cyclohexyl substituents enhanced the Tg of PGA due to rotational barriers within the polymer chains but altered the thermal degradation of PGAs according to TGA. An important issue concerning the production of commercial PGAs is the use of relatively high reaction temperatures for polymerization. In this regard, Schmidt et al. (2014) have recently introduced low-temperature supercritical carbon dioxide as a solvent for the production of PGAs with different molecular masses. This process was not only able to reduce the monomer consumption, heat generation, and the risk of thermal decomposition during polymerization, but it also produced a highly crystalline polymer. The number of average molecular mass (Mn) was found to be independent of thermal parameters and it was also confirmed by changes in the Tg from 31 C to 40 C while Tm was between 217 C and 221 C. These values were found to be in accordance with previous studies for conventionally produced PGA, thus establishing a new synthetic path for biopolymers (Fig. 9.2C and D). More recently, Baez and Marcos-Fernandez (2015) compared the thermal properties of PGAs with similar Mn values but with different alkyl end domains. Higher melting temperatures were observed for the examined PGAs, with a direct relation to the number of methylenes in alkyl end chains.

9.2.2 Copolymers Changing the thermal properties to manipulate biodegradability is still a challenging hurdle for tissue engineering applications. One of the most important approaches to tailor these properties in PGA is the introduction of a new monomer during the polymerization to induce highly biodegradable PGA copolymers. In this regard, Zhou et al. (2004) synthesized PCL
PGA
polyethylene glycol (PEG)
monomethyl-ether random copolymer using bulk polymerization at 160 C. The resultant copolymer had a mass ratio of 10:15:75 for PEG-mono-methyl-ether, GL, and caprolactone (CL), respectively. The final copolymer had a single Tg at 243.8 C while two Tm points were observed at 20.6 C and 48.4 C. These Tg and Tm values are lower than those for PGA homopolymers, which confirms this copolymer as a promising material for tissue engineering applications (Fig. 9.3A). In another study, Ding et al. (2007) produced a copolymer of PGA and aniline pentamer through a low-temperature polycondensation method. They found that this copolymer is not only an electrically conductive biopolymer, but it also has reasonable thermal degradation properties due to the presence of the aniline functional group in its backbone. Marquez et al. (2015) have recently studied the thermal properties of commercially biodegradable suture (Monosyn) containing a triblock copolymer of GL, trimethylene carbonyl, and caproyl units (Marquez et al., 2014, 2015). The study aimed to understand the crystallization behavior of this polymer, employing calorimetric and microscopic methods. They stated that the crystallization of PGL segments is affected by the soft segments due to changes in the secondary nucleation constant and rearrangement of hard blocks over crystal surfaces and heterogenous nucleation and growth of crystals (Marquez et al., 2014, 2015). More recently, Arican and Mert (2015) established an alternative novel thermo-sensitive copolymer

Figure 9.3 (A) DSC heating curve of PCL
PGA
PEG mono-methyl-ether random copolymer with a mass ratio (10:15:75) of PEG monomethyl-ether, GL, and CL, respectively. [The transition temperatures and melting temperatures of polymers were measured by DSC (SEIKO EXSTAR6000) under a flow of nitrogen at a scanning rate of 10 C/min. The measurements were performed in the range 280 C to 80 C under a flow of nitrogen at a scanning rate of 10 C/min. The midpoint of the second run was used for glass transition temperature (Tg) calculation. Figure 4 in original paper.] (B) The TGA curves of PGA
PEG copolymer with different chain lengths (TGAs were performed on a Perkin
Elmer TGA 4000 under nitrogen atmosphere between 20 C and 600 C at a heating rate of 40 C/min. Figure 7 in original paper). Source: From Zhou, S., Xu, J., Yang H., Deng, X., 2004. Macrom. Mater. Eng. 289, 576. Reprinted with permission from Elsevier and (B) From Arican, M.O., Mert, O., 2015. RSC Adv. 5, 71519. Reprinted with permission from the Royal Society of Chemistry.

158

Thermal Analysis of Textiles and Fibers

of PGA
PEG through ring-opening polymerization. According to TGA results, thermal decomposition of the final copolymer is highly related to the portion of PGA and PEG in the chain. In addition, increasing the number of PEG units in the copolymer resulted in the appearance of two melting points in DSC curves (Fig. 9.3B).

9.2.3 Effect of additives on thermal properties Blending of inorganic particles or different polymers with PGA has permitted the production of composite materials with improved physical, mechanical, and biodegradation properties (Linhart et al., 2001; Spearman et al., 2014, 2015; Kuo and Leou, 2006; Park et al., 2006). There are very few publications regarding the use of fillers on thermal properties of PGA. Linhart et al. (2001) developed carbonated apatite/PGA composite for bone replacement. Based on their thermal studies, the melting enthalpy of the composite was 5.3
6.1 kJ/mol while these values were changed to 4.4
6.3 kJ/mol after incubation in serum. Furthermore, DSC studies revealed that the crystallinity of the composite ranged from 45% to 51%. Recently, Spearman et al. (2015) created a PCL
PGA-deoxyribonucleic acid-wrapped single-walled carbon nanotube through electrospinning method. Only one Tg point was observed for both composite fibers with and without the nano-filler, stating the miscibility of final composite fiber. The carbon nanotube was found to improve the Tg from 41.6 C to 50.4 C, which is resulting from a decrease in macromolecular chain mobility and an increase of interfacial interaction within the polymer matrix (Fig. 9.4A). Other attempts to tune the thermal properties of PGA were based on the addition of biodegradable polymers such as poly(lactide-co-glycolide) (PLA-co-GA) reported by Kuo and Leou (2006), chitin (Park et al., 2006), and PCL (Spearman et al., 2014) into PGA in order to generate a polymeric composite with faster biodegradability. Kuo and Leou (2006) evaluated the thermal properties of PGA/(PLA-co-GA) composite scaffold, and they found that the Tg, Tm, and thermal degradation rate were highly dependent on the type of solvent used in the solvent-casting method. In a different study, Park et al. (2006) produced PGA
chitin composite fibers by electrospinning. However, they could not achieve highly miscible scaffolds due to similar melting parameters to the neat PGA (Fig. 9.4B). This result was not observed in the recent report by Spearman et al. (2014) for PCL/PGA composites in which the Tg peaks of PCL and PGA disappeared, indicating a compatible composite with improved mechanical properties.

9.3

Thermal properties of polylactic acid/poly(lactide)s

9.3.1 Homopolymers Polylactic acid (PLA) is a biodegradable linear polyester obtained from renewable sources, such as corn starch, sugar cane, and sugar beets, among others (Spearman et al., 2014; Drumright et al., 2000; Sodergard and Stolt, 2002; Garlotta, 2001).

Figure 9.4 (A) DSC curves of PCL
PGA-blended fibers containing DNA-wrapped single-walled nanotubes during heating (Perkin
Elmer Diamond DSC performed in a liquid nitrogen atmosphere at rates of 10, 20, 30, and 40 C/min from 2100 C to 250 C. Figure 5 in original paper). (B) DSC thermograms of PGA/chitin blend nanofibers prepared by electrospinning with different percentage (%) of each composition (DSC measurements were conducted with a Perkin
Elmer DSC7 instrument in a nitrogen atmosphere. Approximately 10 mg samples was sealed in an aluminum pan for the measurements. The samples were heated from 10 C to 250 C at a rate of 20 C/min, held at 250 C for 1 min, and then quenched to 0 C. The samples were reheated to 250 C. Figure 2 in original paper). Source: From Spearman, S.S., Irin, F., Rivero, I.V., Green, M.J., Abidi, N., 2015. Polymer 56, 476. Reprinted with permission from Elsevier and (B) From Park, K.E., Kang, H.K., Lee, S.J., Min, B.M., Park, W.H., 2006. Biomacromolecules 7, 635. Reprinted with permission from the American Chemical Society.

160

Thermal Analysis of Textiles and Fibers

Having it considered as a suitable replacement for synthetic polymers in certain applications, its use has been considerably hindered due to high production costs (Drumright et al., 2000; Garlotta, 2001). Lactic acid was first isolated in 1780 by Carl Wilhelm Scheele from sour milk, and the ring-opening polymerization of lactide was reported in 1932 by Carothers (Sodergard and Stolt, 2002). These polymers were largely disregarded due to their instability under humid conditions, and it was in the 1960s until their potential in medical applications was first reported (Sodergard and Stolt, 2002). Further studies have shown that PLA provides good grease and oil resistance and is a good barrier to flavors and aromas (Sodergard and Stolt, 2002; Garlotta, 2001). PLA has a high modulus and strength but lacks toughness, which can be amended by orientation, blending, or copolymerization (Drumright et al., 2000; Sodergard and Stolt, 2002; Garlotta, 2001). There are two main pathways to produce PLA: direct condensation of lactic acid, where water removal tends to limit the molecular mass of the final product, and the ring-opening polymerization of lactide, which has been achieved in suspension, solution, melt, and in bulk (Spearman et al., 2014; Drumright et al., 2000; Sodergard and Stolt, 2002; Garlotta, 2001; Lim et al., 2008). Lactic acid is chiral as it has two stereoisomers of L (1) and D (2) (Fig. 9.5) (Sodergard and Stolt, 2002; Garlotta, 2001; Wouwe et al., 2016), which demonstrates how they are affected by polarized light (Lunt, 1998). The amount of L- or D-lactide has a significant role in the crystallinity of the final PLA polymer, which can fluctuate between that of an amorphous polymer to that of a highly crystalline polymer (Petersson et al, 2007). PLAs obtained by the direct condensation polymerization of lactic acid consist mostly of lactyl units, and these polymers often exhibit low-molecular masses since the presence of water can cause chain transfer, resulting in superior mechanical properties (Sodergard and Stolt, 2002; Garlotta, 2001; Lim et al., 2008; Wouwe et al., 2016, Lunt, 1998; Petersson et al., 2007; Maharana et al., 2009, Perego et al., 1996). Lactide is the lactic di-ester of lactic acid, having three enantiomeric forms of D,D-lactide, L,L-lactide, D,L-lactide, or meso-lactide (Wouwe et al., 2016). Among the earlier publications, the effects of the contents of L- and D-enantiomer on the thermal properties have been extensively studied. Fischer et al. (1973) concluded that the melting point of lactide is highly dependent on the content of the

Figure 9.5 Lactic acid (Figure 1.1 in original book). Source: From Groot, W., van Krieken, J., Sliekersl, O., de Vos, S., 2010. In: Auras, R., et al. (Eds.), Poly(Lactic Acid): Synthesis, Structures, Properties, Processing, and Applications. John Wiley & Sons, Inc. (Groot et al., 2010). Reprinted with permission from John Wiley and Sons.

Thermal properties of aliphatic polyesters

161

D-stereo isomeric forms. They also calculated an enthalpy of fusion as 93.6 J/g for PLA with infinite crystal size. Vasanthakumari and Pennings (1983) determined that Tg of PLA is directly proportional to molecular mass of polymer. It was also found that PLAs with different molecular masses represent different crystal morphologies including spherulites at 124 C, coarse grained spherulites at 156 C, axialites at 163 C, and single crystals at 164 C. These crystal forms are induced by processing methods, such as melt or cold crystallization and solution spinning at low drawing temperatures (Mazzullo et al., 1992). In another study, Hyon et al. (1984) performed a melt-spinning process of poly(L-lactide) at 200 C, and the spun fibers were drawn in air at hot temperatures ranging between 80 C and 160 C. Using DSC, the Tg and Tm of fibers were found to be 57 C and 184 C, respectively, by highly oriented fibers. Hoogsteen et al. (1990) produced solution-spun PLA fibers and calculated the melting enthalpy to be in the order of 60
65 J/g. Migliaresi et al. (1991a,b) studied the effects of molecular masses and thermal history on the crystallization of PLAs. Low-molecular mass PLA was crystallized from the melt at cooling rates of 5, 1, and 0.5 C/min, regaining 65%, 79%, and 87% of crystallinity, whereas high-molecular mass PLAs regained their crystallinity under fast cooling rates. These studies reveal that the molecular mass has a profound impact on the crystallinity of PLAs. It has been demonstrated that amorphous PLAs, whether PDLLA or PLLA, exhibit a Tg of around 65 C (Wouwe et al., 2016; Migliaresi et al., 1991a,b). Celli and Scandola (1992) observed that the physical properties of PLAs vary in the range from room temperature to Tg over time, and the polymer eventually becomes brittle due to aging. Tensile and flexural properties of L-enantiomer PLAs are much more affected by its crystallinity than those of D-enantiomer, at which the crystallinity is closely linked to the molecular mass and the stereoregularity of the PLA (Tsuji et al., 2008; Witzke, 1997). Witzke (1997) was also able to correlate the Tg of unoriented poly(L-lactide-co-meso-lactide) to the molecular mass and the contents of L-lactide and meso-lactide. In this study, the Tg at infinite molecular mass for PLLA, poly(meso-lactide), and PDLLA was found to be 61 C, 46 C, and 53 C, respectively. In another study, Dorgan et al. (2005) studied the effects of the L-enantiomer on Tg values of PLAs. It was concluded that Tgs were directly proportional to the contents of L-enantiomers, and they vary between 60.2 C, 56.4 C, and 54.6 C at infinite molecular mass of 100%, 80%, and 50% L-content, respectively (Fig. 9.6), according to Dorgan et al. (2005). The low crystallization rate is one of the most attractive characteristics of PLAs for certain applications, such as production of thermoformed containers or stretchblown bottles; however, the amorphous form has limited applications due to low Tg (Di Lorenzo, 2005; Li and Huneault, 2007). PLA degrades by hydrolysis after prolonged exposure to water; therefore, its biodegradability is significantly studied for tissue engineering applications. Biodegradation of this polymer is mostly caused by chain scission of ester bonds and is greatly influenced by crystallinity, molecular mass, purity, and chain orientation (Wouwe et al., 2016; Ikada and Tsuji, 2000; Pistner et al., 1993; Bergsma et al., 1995).

162

Thermal Analysis of Textiles and Fibers

60

58

Tg/°C 56

100/0 80/20 50/50 0/100

54

52 0

200

400

600

800

10–3Mn

Figure 9.6 Tg values for PLAs with various L-content as a function of number average molecular mass Mn [Perkin
Elmer DSC-7 after calibration with indium. Glass-transition temperatures (Tgs) listed are half-step-height values. Reported melting points are peaks of melting endotherms in second heating scans at 10 C/min after cooling from the melt to room temperature at the same rate. For the homopolymers, the tabulated Tg values refer to specimens quenched as quickly as possible in the instrument from the melts to room temperature to suppress crystallization. Figure 2 in original paper]. Source: From Dorgan, J.R., Janzen, J., Clayton, M.P., 2005. J. Rheol. 49, 607. Reprinted with permission from AIP Publishing LLC.

9.3.2 Copolymers PLAs have been copolymerized with compounds, such as glycolic acid or glycolide, PEG, poly(propylene oxide), (R)-β-butyrolactone, δ-valerolactone (δ-VL), ε-valerolactone, 1,5-dioxepan-2-one (DXO), trimethylene carbonate, and others (Wouwe et al., 2016; Bergsma et al., 1995; Barakat et al., 2001; Khang et al., 2002; Tsuji et al., 2008; Kimura et al., 1989; Stridsberg and Albertsson, 2000; Ryner and Albertsson, 2002; Kurcok et al., 1992; Kesenci et al., 2000). Low-molecular mass PLA-co-GA has been obtained by step-wise polymerization of lactic acid and glycolic acid. The same group produced high-molecular mass PLAco-GA by ring-opening polymerization of lactide (L- or D-) (Wouwe et al., 2016). As the content of the glycolide comonomer increased, the melting point of copolymer decreased as a result of tacticity reduction (Bergsma et al., 1995). Also, the amount of GA monomers accelerated the biodegradation rate of the final copolymer (Barakat et al., 2001). Other study groups observed that degradation times can range from weeks to over a year depending on the fraction of GA (Khang et al., 2002). It is worth noting that small amounts of PGA were found to enhance the crystallinity yield in the copolymer, resulting in higher thermal stability (Tsuji et al., 2008). Block copolymers of PLA and poly(propylene oxide) were synthesized by Kimura et al. (1989), at 150 C using trimethyl aluminum-H2O as a catalyst after

Thermal properties of aliphatic polyesters

163

copolymerization with L-lactide and polypropylene glycol (PPG). This copolymer was melt-spun into fibers resulting in a decrease in Tm, tensile strength, and modulus and an increase in flexibility depending on the PPG content in copolymer. However, there was no considerable change in Tc of the resulting copolymer (Fig. 9.7). Block copolymers of PLA and DXO have been synthesized, aiming to control tensile properties and hydrophilicity (Stridsberg and Albertsson, 2000; Ryner and Albertsson, 2002). It was found that blocks containing PLA have undergone phase separation, resulting in the formation of crystal domains. Thermal characterization of DXO homopolymer showed an amorphous structure with a Tg at 236 C. On the other hand, PLA showed Tg and Tm at 55 C and 169 C, respectively. However, copolymerization of PLA with DXO resulted in the reduction of Tg in comparison with pure PLA, with a Tg value that is closer to the Tg of DXO for triblock and multiblock PLLA
DXO copolymers. Tm values of the resultant copolymers were also decreased to lower values depending on the percentage of DXO monomers in the copolymer structure. AB-type block copolymers of L-lactide and δ-VL were synthesized by anionic polymerization in tetrahydrofurane solution at 20 C, using an alkali-metal alkoxide (potassium methoxide) as initiator (Kurcok et al., 1992). This study was conducted to evaluate the melting enthalpies of block copolymers. An increase in the L-lactide content resulted in an increase in the melting enthalpy at low temperatures, while it showed reverse behavior at higher temperatures. Another

Figure 9.7 Typical DTA curves of PLA homopolymer and block copolymers. Line A is PLA homopolymer and the other lines are copolymers. The ratio of PLA is decreased from the curves of B to F (DTA was recorded under a nitrogen atmosphere on a Shimazu DT-30 thermal analyzer. The rate of heating was set at 10 C/min for a 5.0 mg sample, Figure 2 in original paper). Source: From Kimura, Y., Matsuzaki, Y., Yamane, H., Kitao, T., 1989. Polymer 30, 1342. Reprinted with permission from Elsevier.

164

Thermal Analysis of Textiles and Fibers

interesting result was the immiscibility of L-lactide and δ-VL, which is observed by the generation of two separate melting endotherms in the thermal graphs.

9.3.3 Effect of additives on thermal properties PLA has been considered a suitable binder for particulate bioceramics. Composites of PLA and hydroxyapatite (HA) have been studied for tissue engineering applications, such as hard tissue repair in orthopedics and dental materials. PLA/HA composites were prepared by Kesenci et al. (2000) using solvent casting in chloroform. The resulting composites were amorphous with a greater chain mobility in comparison with PLA. Thermal properties were dependent on HA content: Tg in the range between 54 C and 63 C varies slightly according to the amount of HA, melting points remained between 176 C and 180 C in all compositions; meanwhile, the enthalpy change (ΔHf) is greatly influenced by HA content with variations as large as 43.5 J/g for neat PLA to 19.4 J/g when HA content was 50% mass. Both hydrolytic and thermo-oxidative degradations were observed while studying the thermal properties. TGA results showed complete thermal degradation of composites into vaporizable monomers of lactic acid and oligomers at 350 C, with an increasing degradation rate with increasing HA content in the PLAs with two different molecular mass. In a latter study, Kesenci et al. (2001) reinforced PDLLA/ε-CL blends with silk particles in acetone solution. Before blending, the copolymer had a 50/50 PDLLA/ε-CL ratio with a Tg around 219.9 C. After the solvent was evaporated, transparency of the composites decreased as the filler content increased, a homogenous filler distribution was observed, and the Tg values of the composites were 218.9 C (10% filler), 214 C (30% filler), and 217.2 C (50% filler). Onset degradation temperature was 160 C for 50/50 filler/copolymer composite while all the other composites degraded above 230 C. Several researchers used clay nanoparticles as filler with PLAs and studied the performance of the resultant composites (Ray et al, 2003; Lee et al., 2003a,b). PLAs have been blended with montmorillonite (MMT) that is modified with octadecylammonium cation by melt extrusion (Ray et al., 2003). The silicate layer was intercalated by PLA and randomly distributed within the polymer. In one study, oligo(ε-CL) was used to improve the compatibility of the composite and to enhance the stacking of the clay layers. The resulting composites containing 4, 5, and 7 mass% modified MMT showed better mechanical properties than neat PLA. Based on thermal analysis, the Tg and Tm of the composites were marginally modified, in which a significant improvement of biodegradability was observed, and the composites were completely degraded after 2 months. Enhanced biodegradation of PLA
MMT composites was also observed in another study by Lee et al. (2003a,b). MMT nano-platelets modified with dimethyl dehydrogenated tallow ammonium cations were used for composite production. Lower Tg values were observed for the produced composites because the mobility of PLA backbones was increased, which eventually resulted in a higher amorphous region of the matrix (Fig. 9.8A).

Figure 9.8 (A) DSC curves of quenched specimens of PLA/MMT composite containing different amounts of MMT measured at 10 C/min. (TA Instruments DSC 2910 at a heating rate of 10 C/min in nitrogen environment. Figure 2 in original paper.) (B) TGA analysis of PLA plasticized with PEG 1000 and filled with 3 mass% of (organo-modified)-montmorillonite. Cl. stands for Cloisite and p. PLA for plasticized PLA [TGA was performed using a Hi-Res TGA 2950 thermogravimetric analyzer with a heating ramp of 20 K/min under air flow (74 cm3/min) from room temperature to 600 C. Figure 3 in original paper]. Source: From Lee, J.H., Park, T.G., Park, H.S., Lee, D.S., Lee, Y.K., Yoon, S.C., et al., 2003a. Biomaterials 24, 2772. Reprinted with permission from John Wiley and Sons and (B) From Paul, M.A., Alexandre, M., Dege´e, P., Henrist, C., Rulmont, A., Dubois, P., 2003. Polymer 44, 443. Reprinted with permission from the Royal Society of Chemistry.

166

Thermal Analysis of Textiles and Fibers

PLLA was blended with nanosized MMT (both alkyl ammonium cationmodified and unmodified) using PEG as a plasticizer by direct melt blending (Paul et al., 2003). The effect of the clay on Tg and Tm of composites was negligible, while the thermal stability of the composites was improved depending on the nature of the alkyl-ammonium cations (length of the alkyl chain or functionality attached to the ammonium cation) (Fig. 9.8B). A similar study was conducted by Wu and Wu (2006) by using modified MMT as filler into PLA in solution. MMT was modified with n-hexadecyl trimethylammonium bromide (CTAB) cations and further modified with chitosan in an aqueous solution with 1 mass% lactic acid. They found that the crystallinity and ΔHm decrease from 23% and 19 J/g to 14% and 2 J/g, respectively, as the MMT content increases from 1.5 to 6 mass% in composites. Recent studies were conducted to produce PLA composites with cellulose-based particles and carbohydrates (Shakoor et al., 2013; Herrera et al., 2016). Shakoor et al. (2013) reinforced PLA with hemp fibers at 10
30 mass%. The Tg of the composite with the highest filler content was increased in comparison with pure PLA due to stabilizing capability of the hemp fibers. More recently, PLA was reinforced with cellulose and chitin nanocrystals by extrusion, using triethyl citrate as plasticizer and ethanol as solvent (Kulizinski and Piorkowska, 2005). It was demonstrated that faster cooling rates produced composites with better elongation at break and transparency, and the composites that are reinforced with chitin nanocrystals offered better mechanical properties than those reinforced with cellulose (Herrera et al., 2016). Cellulose significantly increased the nucleation of spherulites, and those composites exhibited a slightly higher crystallinity despite having a lower melt crystallization temperature. Plasticizers and nucleating agents have shown an effect on thermal properties and crystallization rate of PLAs. PEG was used as plasticizer to improve the processability of PLA (Kulizinski and Piorkowska, 2005). However, the cold crystallization temperature and Tg of PLA were decreased due to better chain mobility and higher spherulite growth rate. Talc has also been shown as an effective nucleating agent in combination with plasticizers, such as PEG and acetyl triethyl citrate, resulting in an increase in the crystallization percentage under high cooling rates (Kulizinski and Piorkowska, 2005). Kawamoto et al. (2007) evaluated the effect of several methylene chain numbers in hydrazide compounds as nucleating agents on thermal properties of PLA. A benzoyl hydrazide compound was the most effective nucleating agent in this study to improve the PLA crystallization. In contrast, decamethylenedicarboxylic dibenzoylhydrazide provided the highest crystallization temperature and enthalpy (131 C and 46 J/g, respectively) in PLA, making it comparable to nucleated poly(propylene).

9.4

Thermal properties of polycaprolactones

9.4.1 Homopolymers Poly-ε-caprolactone (PCL) is a hydrophobic, semi-crystalline, bioresorbable aliphatic polyester derived from the relatively inexpensive monomeric unit of ε-CL to

Thermal properties of aliphatic polyesters

167

Figure 9.9 The chemical structure of PCL biopolymer. PCL, Polycaprolactone.

give linear hexanoate repeating units (Fig. 9.9). Its biodegradability, biocompatibility, and nontoxic properties attracted significant attention in biomedical studies related to sustained drug delivery, artificial skins, and artificial bones (Pitt, 1990; Sinha et al., 2004; Huang et al., 2003; Woodruff and Hutmacher, 2010; Coombes et al., 2004; Pitt et al., 1981). Having a degree of crystallinity that could reach up to 69%, PCL is found to exhibit decreasing crystallinity with increasing molecular mass (Yu et al., 2015; Pitt, 1990; Mark, 1999; da Silva et al., 2011). Synthesis of the PCL polymer is found to follow two main pathways: either through the ringopening polymerization of ε-CL using a combination of catalysts which are anionic, cationic, and coordination based or through the ring-opening polymerization of 2-methylene-1-3-dioxepane (Yu et al., 2015; Pitt, 1990). PCL exhibits tailorable degradation kinetics related to thermal and mechanical characteristics, both of which primarily depend on its molecular mass, degree of crystallinity, and conditions of degradation (Mark, 1999). The low Tm characteristic of PCL typically ranges between 56 C and 67 C and is related to its high flexibility and elongation at break (Hatakeyama et al., 2000; Balsamo et al., 2006; Albuerne et al., 2003; Arnal et al., 2004). Elastic properties of PCL are attractive in blends with other polymers to either enhance biocompatibility or modulate mechanical properties for tissue engineering applications (Labet and Thielemans, 2009; Matzinos et al., 2002; Hatakeyama et al., 2000). With an average molecular mass ranging between 530,000 and 630,000 g/mol for most PCL grades, the Tg of PCL ranges between 265 C and 260 C, which is low enough to be rubbery at room temperature (Van de Velde and Kiekens, 2002; Schlesinger et al., 2015; Song et al., 2016; Wang et al., 2016; Ikada and Tsuji, 2000). The Tc of PCL was widely reported in the literature to be around 30 C (Balsamo et al., 2006; Albuerne et al., 2003; Arnal et al., 2004). At the Tm of 67 C, the melting enthalpy (ΔHf) of 66% crystallized PCL was reported to be 88 J/g, while ΔHf of 100% crystalline PCL is 139.5 J/g. Hatakeyama et al. (2000) have found a trend in the rearrangement of crystalline regions, establishing a direct relation with molar ratio of CL to hydroxyl groups. Many studies have addressed the effect of different processing parameters and conditions on the Tg and Tm characteristics of PCL. Sravanthi et al. (2010) developed different porous structures of PCL scaffolds and studied the melting endotherm of the products. The Tms were found to be in the range of 60.5 C
63 C, close to the melting endotherm of the as-received PCL. da Silva et al. (2011) studied the melting characteristics of iodine polymerized PCL in comparison to PCL received from Sigma-Aldrich. They observed a broad melting temperature for produced PCL with the calculated Χc of 78%, whereas it was 51% for commercially available PCL. Several studies have been done to evaluate the thermal behavior of oxidized and gamma-irradiated PCLs in comparison to the neat PCL. Results indicated lower Tg,

168

Thermal Analysis of Textiles and Fibers

(i) and (iii) PCL (iii) and (iv) PCL-OX

Heat flow (ENDO)

(i) (ii) (iii)

5 mW

(iv)

–20

0

20

40

60

80

100

Temperature (°C)

Figure 9.10 DSC curves for neat PCL and oxidized PCL. (Thermal analysis was performed by standard DSC and by applying SSA technique. DSC and SSA were performed with a Perkin
Elmer DSC-7 under an ultrahigh-purity nitrogen atmosphere. The equipment was calibrated with cyclohexane and indium standards. Sample weight was kept constant at 5.0 mg. In order to perform a preliminary evaluation of the thermal behavior of the sample, standard heating and cooling scans were recorded from 220 to 130 C at 10 C/min. All melting parameters were determined from second heating scan. Figure 4 in original paper.) Source: From Sabino, M.A., 2007. Polym. Degrad. Stab. 92, 986. Reprinted with permission from Elsevier.

Tm, and Tc values as well as smaller crystals for modified PCLs due to generation of new functional groups and PCL chain scission (Fig. 9.10) (Balsamo et al., 2006; Arnal et al., 2004; Sabino, 2007; Brito et al., 2006; Rojas de Gascue et al., 2002; Mu¨ller and Arnal, 2005; Sabino et al., 2001; Darwis et al., 1998). Such thermal behaviors of modified PCLs are important for any approaches toward producing highly hydrolytic and biodegradable polymers. Regarding thermal degradation of PCL, several studies have proposed that it occurs in two successive steps: first, via pyrolysis at the ester linkage of PCL, causing a propagated rupture of the polyester chains, followed by an unzipping depolymerization and release of ε-CL monomers (Persenaire et al., 2001; Aoyagi et al., 2002). Release of ε-CL has been observed in many studies upon the hydroxyl chainend scission of PCL during pyrolysis and supercritical degradation (Sivalingam et al, 2003a; Joshi and Madras, 2008; Yeo and Kiran, 2005). The mode of scission of the PCL chains is highly dependent on temperature at which thermal degradation is occurring. It is noteworthy that it starts with end-chain scission at higher temperatures (supercritical) and random-chain scission at lower temperatures (subcritical) (Yeo and Kiran, 2005; Goto et al., 2006; Becker et al., 2015). Thermal degradation of PCL in solution and in vivo has attracted significant attention due to the direct applications in sustained drug release, biomedical device implants, surface coatings, and adhesives (Sinha et al., 2004; Woodruff and Hutmacher, 2010; Yeo and Kiran, 2005).

Thermal properties of aliphatic polyesters

169

The thermal degradation of PCL in bulk was compared to that in solution and it was found that PCL degrades by random-chain scission in bulk while it degrades by specific chain-end scission in solution (Sivalingam and Madras, 2003; Sivalingam et al., 2004; Dhawan et al., 1991). Joshi and Madras (2008) studied the thermal degradation of PCL in both supercritical and subcritical solvents of toluene, ethylbenzene, o-xylene, and benzene with a temperature range of 250 C
300 C (subcritical) and 300 C
375 C (supercritical). In toluene, PCL was found to degrade by randomchain scission in subcritical conditions and by chain-end scission in supercritical conditions, while it showed a chain-end scission behavior in both subcritical and supercritical conditions for other solvents. This observation demonstrated that the mode of scission is not directly related to the criticality of the solution but rather to the temperature of the solution, with a decreasing rate of degradation at elevated pressures (Yeo and Kiran, 2005; Goto et al., 2006). It is well known, however, that polymers tend to undergo accelerated degradation rates when in supercritical solutions. Dhawan et al. (1991) described that one way to enhance the yield and selectivity of pyrolysis reactions is to use proper solvents at supercritical conditions as reaction media, which in turn decreases the operating temperature of the reaction. The relation between biodegradation and thermal stability of PCL has been thoroughly studied due to the promising characteristics of PCL and its critical applications (Sivalingam et al., 2003a,b; Kobayashi et al., 2000; Brode and Koleske, 1972; Megna and Koroscil, 1968). Several parameters influence the thermal and hydrolytic stability of polymers, such as the distinct type of chain-end groups, the catalysts used, and their respective concentrations. Brode and Koleske (1972) developed hydrolytic and thermally unstable linkages in PCL by incorporation of lactone into the tin hetero-atom bond of stannous octanoate and trimethyl tin acetate, and also by preparation of PCL using phosphorus pentafluoride. Despite these studies, tetrabutyl titanate and other metallic catalysts can be used at a temperature range between 250 C and 300 C to depolymerize PCL into its monomeric form; dimers, trimers, and other oligomeric products in smaller proportions (Chen et al., 2000). Degradation studies on PCL have addressed both enzymatic and hydrolytic degradation (Woodruff and Hutmacher, 2010; Pitt et al., 1981). Although PCL degradation can be catalyzed by enzymes in the environment which in turn can depolymerize PCL at a faster rate, PCL does not exhibit enzymatic degradation in the body (Ikada and Tsuji, 2000; Gan et al., 1997). Several studies analyzed in vitro versus in vivo enzymatic degradation of PCL, and it was found that it did not exhibit a significant role in the elementary phases of degradation (Sun et al., 2006; Ali et al., 1994; Yoon and Ji, 2005; Ouhadi et al., 1976; Kwon et al., 2001; Wang et al., 2000; Lee et al., 2003a,b; Min et al., 2005). Like other heterochain polymers, PCL is susceptible to hydrolytic degradation under acidic or alkaline conditions with a significantly accelerated degradation rate at elevated temperatures. Similar to many other polymers, the conditions in the human body are aggressive enough to stimulate hydrolytic degradation of PCL at 37 C (Yoon and Ji, 2005). Thermal degradation of PCL is found to follow first-order kinetics with a degradation rate equal to in vitro hydrolysis at 40 C. Yoon and Ji (2005) stated that PCL loses half of its

170

Thermal Analysis of Textiles and Fibers

tensile strength within 2 months of in vitro treatment. On the contrary, oxidative degradation occurred in PCL upon oxygen contact at 120 C (Kwon et al., 2001).

9.4.2 Copolymers The majority of PCL copolymers are produced with PGA, polylactide (PLA), PEG, and polyurethane (PU) (Wang et al., 2000; Lee et al., 2003a,b; Min et al., 2005, Jia et al., 2008; Ugartemendia et al., 2014; Maglio et al., 2003; Hoque et al., 2009; Bogdanov et al., 1998; Van Bogart et al., 1983; Seefried et al., 1975; Kweon et al., 2000). Copolymers of PCL often have a lower melting temperature and lower crystallization rate, to make them more prone to degradation. A key copolymer which has been thoroughly studied in literature is poly(glycolide-co-caprolactone) (PGCL) which consists of the strong and rigid PGA and the soft and flexible PCL segments. In a study by Kwon et al (2001), PGCL random copolymer PGA:PCL (51:49) was found to be relatively elastic with a superior elongation at break compared to PGA: PCL (70:30). The Tg of PGA:PCL (51:49) was 215.4 C, while no Tc and Tm points were reported, suggesting a random and amorphous structure in the resultant copolymer. In another study, Wang et al. (2000) synthesized a segmented block copolymer of PGA:PCL (75:25) in order to determine the effect of composition on thermal properties and crystallinity in comparison to PGA homopolymer. Another part of their study was to compare the thermal properties of different random copolymers of PGA:PCL (5:95, 90:10). They reported the Tms of PGA:PCL segmented block copolymers as 217 C, while the Tms of PGA:PLA (5:95) and PGA:PLA (90:10) were 173 C and 201 C, respectively. It was also found that the degree of crystallinity and molecular orientation dominated the degradation behavior. In addition, PGA has a faster isothermal crystallization rate compared to PCL; therefore, PGA:PCL (75:25) segmented block copolymer had the highest crystallization rate compared to neat PCL (Lee et al., 2003a,b). DSC also confirmed the lowest crystallinity degree of PGA
PCL (75:25) copolymer compared to PGA:PLA (5:95) and PGA:PLA (90:10). Lower degree of crystallinity was also obtained for PGA:PCL copolymer by ring-opening polymerization of GL and CL in the presence of tin octanoate catalyst. They were able to achieve an amorphous PGA:PCL copolymer with a Tg at 219.3 C, while having no Tc and Tm in the curves. Since biopolymers used in tissue engineering are required to exhibit reasonable biodegradation rates and low Tg, Min et al. (2005) obtained PLA
PGCL multiblock copolymer by ring-opening polymerization and coupling reaction. They were not only able to produce a copolymer with a higher degradation rate and lower Tm than those in PCL but also achieved a highly cross linked copolymer with shape memory properties. It is noteworthy that Tm values of this multiblock copolymer were significantly influenced by the molecular mass of the PLA and PGCL macrodiols. They also reported that the PLA
PGCL copolymer exhibited a Tg and a Tm above 45 C and 120 C, respectively. Lactide monomer has also been copolymerized with CL to alter PCL thermal properties. Jia et al. produced a series of PCL
trimethylene carbonate
PLA terpolymers and evaluated their thermal properties by DSC method. Thermal characterization showed only one Tg at

Thermal properties of aliphatic polyesters

171

14.7 C for the terpolymer with the ratios of 40:10:50, respectively, for PCL, trimethylene carbonate, and PLA blocks. It was observed that increasing the trimethylene carbonate content to 50% reduced the Tg to 232.7 C in the terpolymer. This observation indicated that a continuous amorphous phase was present in the terpolymer, due to the random placement of trimethylene carbonate in the chain. On the other hand, Tm for the terpolymer was near the Tm of PLA homopolymer and decreased from 133.8 C to 118.2 C on decreasing the lactide content from 50 to 40 (molar ratio) (Jia et al., 2008). Ugartemendia et al. (2014) have recently synthesized PLA
PCL copolymers and studied the thermal properties after quenching. They found that the copolymer has highly amorphous morphology with a Tg at 25.5 C in contrast to PLAs that are semi-crystalline. PEG can be copolymerized with PCL due to its water solubility and subsequent compatibility with cells (Hoque et al., 2009). Thus, PEG/PCL copolymers were studied with the aim of linking thermal properties to composition. The melting temperatures of di- and triblock PEG
PCL copolymers were found in the temperature range of 63 C
66 C, similar to PCL homopolymer. However, the presence of PEG decreased the water contact angle and resulted in the improvement in the cell culture performance of the final copolymer (Bogdanov et al., 1998; Van Bogart et al., 1983). Diblock, triblock, and star-copolymers were synthesized by Bogdanov et al. (1998), they studied the effect of PEG
PCL ratio on the thermal properties of the resultant copolymers. It was observed that the PCL segments crystallized first when cooling from the molten state, leading to incomplete crystallization of the PEG blocks. The thermal properties and morphology of copolymers were found to be dependent on the chain length of the PCL and PEG constituents (the sequence of different blocks). Accordingly, all copolymers showed two distinct exothermic peaks at 30 C
33 C and 214 C corresponding to the crystallization temperatures of PCL and PEG blocks, respectively, which confirmed the dependency of thermal properties on molecular mass of each polymer block. In addition, the crystallization temperature of the PCL block (first exotherm) decreased with decreasing molecular mass of the PCL block. Another observation was appearance of two Tm peaks at 55 C
58 C and 25 C
50 C for the copolymer attributing to melting of PCL and PEG crystals, respectively. They also demonstrated that the PCL blocks crystallize first and change the arrangement of the copolymer structure leading to imperfect crystallization of the PEG blocks. In a study on the PCL
PU copolymer by Van Bogart et al (1983), thermal properties of copolymer were studied based on the changes in the hard-segment content (23%
77% by mass). The size and peak of the Tm, as well as the crystallinity percent of copolymer, were increased in relation to the increasing of hard-segment content at constant soft segment molecular mass (830 or 2000 MW PCL). In a similar study, a series of thermoplastic urethane elastomers was produced by varying the sequence lengths of PCL diols and 4,40 -diphenylmethane diisocyanate and 1,4-butanediol (Najemi et al., 2010). This study has shown a direct dependency of the Tg on the sequence length of the polymer soft segments, where a trend of decreasing Tg was observed with increasing softsegment chain lengths. Although PGA, PLA, PEG, and PU were mostly used for copolymerization with PCL, some studies have attempted the PCL blending of

172

Thermal Analysis of Textiles and Fibers

macromonomers with starch (Jayakumar and Tamura, 2008; Herweh et al., 1977), chitin (Lovinger et al., 1993), poly[(1,4-cyclohexane) bismethylene isophthalate] (Xiong et al., 2006), poly(dimethyl siloxane) (Wu et al., 2014), polyethylene oxide
polypropylene oxide blocks (Lin, 1999), poly[(N-vinylcaprolactam)] (Trujillo et al., 2012), δ-VL, γ-butyrolactone, γ-valerolactone, and γ-CL (Chin et al., 2015). Table 9.1 reports the thermal parameters of different PCL copolymers according to previous studies.

9.4.3 Effect of additives on thermal properties An alternative method in the application of biocomposites in tissue engineering is developing PCL composites with nanoparticles or polymers, which in turn affects the thermal and mechanical properties. Nanoparticles with a variety of shapes were used in PCL nanocomposite formation including nanotubes and nanospheres (Zhang et al., 2013; Yeh et al., 2009; Dziadek et al., 2015a,b; Wu, 2004; Cesur et al., 2015; Plazas Bonilla et al., 2014; Demirdo¨gen et al., 2014; Lee et al., 2008; Chandra and Rustgi, 1998; Cipitria et al., 2011; Dash and Konkimalla, 2012; Luetzow et al., 2007; Sisson et al., 2013; Tokiwa and Calabia, 2007; Macrae and Table 9.1 Thermal characteristics of PCL copolymers reported by researchers. Copolymer

Tg ( C)

Tm ( C)

References

PGCL (51:49) PGCL (5:95) PGCL (90:10) PGCL (75:25) PGCL PLA
PGCL

2 15.4


2 19.3 45


173 201 217
130
160

PCL
trimethylene carbonate
PLA

2 32.7 to 14.7 depending on segment ratio 25.5
60
62 195 or 238 depending on synthesis method 2 56


133.8
118.2 depending on segment ratio
66
63



Park et al. (2006) Spearman et al. (2014) Spearman et al. (2014) Spearman et al. (2014) Drumright et al. (2000) Sodergard and Stolt (2002) Garlotta (2001)

PLA
PCL PEG
PCL PEG
PCL PCL-starch

PCL-starch PCL
polyethylene oxide
polypropylene oxide


41
31 depending on segment ratio

Lim et al. (2008) Wouwe et al. (2016) Petersson et al. (2007) Vasanthakumari and Pennings (1983)

Mazzullo et al. (1992) Migliaresi et al. (1991b)

Thermal properties of aliphatic polyesters

173

Wilkinson, 1958; Wallen and Rohwedder, 1974). For example, PCL/multi-walled carbon nanotube (MWCNT) nanocomposites were made through melt blending by adding MWCNTs from 0.3% to 5% of the polymer mass. It was observed that MWCNTs are not only able to produce an electrically conductive composite but also increased the Tg and Tc values about 3 C and 14 C, respectively (Fig. 9.11) (Zhang et al., 2013; Yeh et al., 2009). It has been demonstrated that functional nanoparticles play a key role in thermal and mechanical properties of polymer nanocomposites. Due to the hydrophobic nature of PCL, several researchers utilized acrylic acid-grafted PCL for production of PCL nanocomposites containing fillers (Dziadek et al., 2015a). Consequently, the interfacial interactions and bonding between PCL and nanoparticles were increased. Such interactions were more pronounced between OH-grafted carbon nanotubes on one hand and acrylic acid-grafted PCL on the other hand which intensified the increment of Tg, Tc, and Tm of the resultant PCL/MWCNT nanocomposites. Also, further increments can be seen in crystallinity percentage and thermal stability of PCL/MWCNT nanocomposites. This observation was due to the reduction in PCL chain mobility in the amorphous regions occurring by interactions between the filler and PCL (Dziadek et al., 2015b). Bioactive glasses for bone tissue engineering were also examined. These fillers containing different values of

Figure 9.11 DSC (A) second heating and (B) cooling curves for unfilled PCL and PCL/ MWCNT composites (Perkin
Elmer Diamond DSC at a ramp rate of 250 K/min. Samples were heated from 2100 C to 120 C and held at 120 C for 1 min. The samples were then cooled to 2100 C, held at 2100 C for 1 min before being reheated to 120 C at 250 K/min. At least three specimens were tested for each sample and the corresponding mean values reported. Figure 2 in original paper). Source: From Chin, S.J., Vempati, S., Dawson, P., Knite, M., Linarts, A., Ozols, K., et al., 2015. Polymer 58, 209. Reprinted with permission from John Wiley and Sons.

174

Thermal Analysis of Textiles and Fibers

silica and calcium oxide were incorporated into PCL (Wu, 2004; Cesur et al., 2015). In contrast to carbon nanotubes, they are capable of reducing Tc, crystal size, and degree of crystallinity. The nanocomposites of PCL/bioactive glass microparticles were cited for their high hydrolytic degradation rate, bioactivity, and biocompatibility to tissue and cell culturing (Wu, 2004). Similar to carbon nanotubes, silica and titanium dioxide nanoparticles can also improve the thermal parameters of polymers. Several research groups studied the incorporation of SiO2, TiO2, ZnO, and clay nanoparticles in PCL, aiming to improve thermal and mechanical properties (Dziadek et al., 2015a; Plazas Bonilla et al., 2014; Demirdo¨gen et al., 2014; Lee et al., 2008; Chandra and Rustgi, 1998). It was interestingly found that these nanoparticles increase the Tg of PCL due to hindering the long-chain motions in PCL polymer. This result was observed in the PCL nanocomposite while increasing the fraction of nanoparticles up to a certain value. However, any excess in nanoparticle fraction above this threshold has a complicated effect on thermal characteristics due to phase separation between the nanoparticles and PCL matrix. The effect of the addition of organic additives or polymers into PCL on thermal properties of the resultant blends has been a topic receiving considerable interest in the recent years (Cipitria et al., 2011; Dash and Konkimalla, 2012; Luetzow et al., 2007; Sisson et al., 2013). Cesur et al. (2015) have recently studied the effect of oleic acid and glycerol on crystallization properties of the blend with DSC. They found that oleic acid and glycerol reduce the crystallization percent of PCL owing to plasticizing effect of these additives. Other recent studies were performed to evaluate the effects of high-molecular mass polymers, such as polyhedral oligomeric silsesquioxane (POSS), on physical and mechanical properties of PCL. POSS/PCL matrix has shown a high crystallinity (74%) with an increased Tg, Tm, and heat of fusion (Cipitria et al., 2011).

9.5

Thermal properties of polyhydroxyalkanoates

9.5.1 Homopolymers PHAs (Fig. 9.12) were discovered in the 1920s as linear polyesters utilized by bacterial fermentation. It was found that an aerobic bacillus (Bacillus megaterium) was able to produce β-hydroxybutyric acid in anaerobic conditions as a carbon and energy source (Macrae and Wilkinson, 1958). Lemoigne extracted a substance

Figure 9.12 The chemical structures of the most common PHA biopolymers.

Thermal properties of aliphatic polyesters

175

using chloroform in 1944, which he was able to identify as a polymer of β-hydroxybutyric acid. In collaboration with Delaporte and Croson, Lemoigne discovered that the polymer was the major component of “lipid” cytoplasmatic granules that were recovered from the bacillus. Further research suggested that the polymer acts as reserve of carbon and energy in the bacillus. The same phenomena were observed afterward in Gram-negative bacteria (Wallen and Rohwedder, 1974; Sudesh et al., 2000), although it was originally inferred that poly(3-hydroxybutyrate) (P3HB) was the only PHA to be found in the carbon reserve. In the 1970s, polyesters were isolated from activated sludges using primarily chloroform as solvent (Wallen and Rohwedder, 1974; Sudesh et al., 2000). These polyesters, though, similar in some physical and thermal properties to P3HB, were identified as poly(3-hydroxyvalerate) and poly(3-hydroxyhexanoate) (Macrae and Wilkinson, 1958). In the following decade, the presence of up to 11 other HA units was discovered, and in recent years, more than a hundred PHAs have been identified (Sudesh et al., 2000). P3HB is still the most economically feasible PHA to produce, and hence the most abundant. For that reason, most of the studies have been performed on the enhancing of the properties of P3HB homopolymer. Regarding the thermal properties of homopolymers of PHA, most of the data available are about P3HB. Neat P3HB has a Tg of 4 C, a melting point of 177 C, and ΔHf 5 20.1 cal/g (584.1 J/g) (Kunioka et al., 1989). Crystallinity is one of the most influential parameters in the structure, thermal, and mechanical properties of PHAs. The crystallinity of P3HB is well known as the main cause of its stiffness and brittleness, and different research groups have found several values for the crystallinity of neat P3HB: Kunioka et al. (1989) 55% 6 5%, Doi et al. (1990) 60% 6 5%, and 80% Luzier (1992). P3HB brittleness has been attributed to cracks along the spherulites (de Koning and Lemstra, 1993), as the polymer is strained, the cracks get bigger and bigger and coalesce with each other. de Koning and Lemstra (1993, 1994) demonstrated that stiffness and brittleness are caused by progressive crystallization; however, this process can be remediated to some degree by annealing. All of the earlier data were obtained from PHAs extracted from Alcaligenes eutrophus, but Saito and Doi (1994) studied the properties of different copolymers of P3HB and P4HB obtained from Comamonas acidovorans. They obtained the following results for neat P4HB: Tm 5 53 C, Tg 5 248 C, ΔHf 5 8.6 cal/g, and the crystallinity 5 34% 6 5%.

9.5.2 Copolymers Designing PHA copolymers by graft copolymerization with different monomers of the same family or other biopolymers has been extensively studied in the literature (Saito and Doi, 1994; Doi et al., 1995; Hazer and Steinbu¨chel, 2007; Arslan et al., 2007; Hazer, 2015). It is aimed to (1) producing amphiphilic copolymers by adding hydrophilic segments into the hydrophobic PHAs, (2) producing more flexible or rubber-like copolymers, (3) increasing the hydrophilicity of PHAs for certain medical applications, and (4) producing new copolymers with improved thermal,

176

Thermal Analysis of Textiles and Fibers

crystallization, and mechanical properties. Such modifications are necessary for PHAs to improve their diversification in different industries (Saito and Doi, 1994). In this regard, Hazer and Steinbu¨chel (2007) have reported a valuable review on different chemical modification methods of PHAs including grafting reactions, copolymerization, chlorination, cross linking, epoxidation, and hydroxyl and carboxylic acid functionalization. Saito and Doi (1994) synthesized a random copolymer of 3-hydroxybutyric acid (3HB) and 3-hydroxyhexanoic acid (3HH) using olive oil as carbon source. They interestingly found that by increasing the 3HH fraction in copolymerization results in producing a soft and flexible copolymer with lower thermal parameters. DSC results showed decreasing of Tg and Tm from 4 C and 177 C for P3HB copolymer to 24 C and 52 C, respectively, for P(3HB-co-3HH) copolymer with 25 mol% of 3HH fraction. Arslan et al. (2007) introduced atom transfer radical polymerization for synthesis of P3HB-g-poly(methylmethacrylate) brush-type graft copolymers. It was aimed to improve the chemical properties of PHA. However, DSC thermograms of the final P3HB-g-PMMA brush-type graft copolymers exhibited different Tg corresponding to each segment, confirming microphase-separated segments in both block and graft copolymers. This group has recently reported synthesis of amphiphilic poly(3-hydroxy alkanoate)s with pendant hydroxyl and carboxylic groups via photo-click reactions (Fig. 9.13A). This research showed that PHA derivatives with different carboxyl and hydroxyl functionalization cause to change in Tg values from 225 C to 238 C. Similar changes were observed in Tm values of modified copolymers as 45 C, 58 C, and 98 C stating the highest Tm value for hydroxyl-functionalized copolymer (Fig. 9.13B) (Arslan et al., 2007). Researchers have been trying to improve the thermal properties of poly(3-hydroxyoctanoate) (PHO) by copolymerization with glycerols, such as 1,3-diglycerol diacrylate (GDD-g-PHO) using homogeneous solutions of PHO, GDD monomer, and benzoylperoxide initiator under heating. PHO showed the lowest Tg, Tm, and ΔHm, The Tg, Tm, and ΔHf of the graft copolymers increased in comparison with untreated PHO as the DG of the GDD groups in the copolymers increased. This improvement was due to intermolecular and intramolecular hydrogen bonding between GDD-grafted chains that are formed (Fig. 9.13C) (Hazer, 2015; Kim et al., 2008).

9.5.3 Effect of additives on thermal properties Even though PHAs can be modified by copolymerization process, another approach to modify the properties is addition of organic or inorganic components into the polymer. Ke et al. (2017) have recently focused on different additives used during the blending process with PHAs as a review paper. It has been shown that additives can be blended with PHAs as monomers, particles, cross linkers, or other polymers (Ke et al., 2017). Cross linking of PHAs based on peroxidized compounds has been recently considered as an effective method to improve the thermal stability and

Thermal properties of aliphatic polyesters

177

Figure 9.13 (A) Schematic procedure of the carboxylated(a) and hydroxylated(b) PHA derivatives reported by Hazer (2015) (Scheme 2 in original paper). (B) DSC curves of the PHA derivatives: (a) poly(3-hydroxy octanoate-coundecenoate) containing COOH groups, (b) poly(3-hydroxy octanoate-coundecenoate) containing OH groups, (c) poly(3-hydroxy undecenoate) containing COOH groups, (d) poly(3-hydroxy undecenoate) containing OH groups, (e) poly(3-hydroxy octanoate-co-soybean oil polymer) containing OH groups, (f) poly (3-hydroxy octanoate-co-soybean oil polymer) containing COOH groups. [Thermal analysis of the obtained polymers was carried out under nitrogen using a TAQ2000 DSC and Q600 Simultaneous DSC-TGA (SDT) series thermal analysis systems. DSC measures temperatures and heat flows associated with thermal transitions in the polymer samples obtained. The dried sample was heated from 260 C to 120 C under nitrogen atmosphere heating from 20 C to 600 C at a rate of 10 C/min. Figure 7 in original paper.] (C) Schematic diagram reported by Kim et al. (2008) for graft copolymerization of glycerol 1,3-diglycerolate diacrylate onto PHO (Figure 1 in original paper). Source: (A) From Hazer, B. (2015) Polym. Degrad. Stab. 119, 159. Reprinted with permission from Elsevier, (B) From Hazer, B., 2015. Polym. Degrad. Stab. 119, 159. Reprinted with permission from Elsevier; and (C) From Kim, H.W., Chung, M.G., Kim, Y.B., and Rhee Y.H. (2008) Int. J. Biol. Macromol. 43, 307. Reprinted with permission from Elsevier.

178

Thermal Analysis of Textiles and Fibers

mechanical properties of the resultant composite material. In this regard, several researchers used di-cumyl peroxide (DCP) in an extruder by melt blending with PHAs at 160 C (Fei et al., 2004a; Vogel et al., 2007; Ma et al., 2012; Wei and McDonald, 2015). Fei et al. (2004a) indicated that the Tm and Tc of Poly(3-hydroxybutyrate-co-3-hydroxyvalerate)(PHBV) decreased significantly with increasing DCP content (Fig. 9.14A). However, Vogel et al. (2007) found that DCP does not affect the Tg of P3HB while it changes significantly the shape of Tm peaks in the blended composites. In more recent years, researchers have been using DCP as cross linker between PHAs and other polymers, such as poly(butylene succinate) (Vogel et al., 2007), PLA (Ma et al., 2012), and triallyl trimesate (Wei and McDonald, 2015). Ma et al. (2012) prepared Poly(3-hydroxybutyrate-co-3-hydroxyvalerate)/Polybutylene succinate(PHBV/PBS) blends using DCP by melt blending and observed improvement in thermal stability of composite materials in comparison with the blank PHBV (Fig. 9.14B). Wei and McDonald (2015) have recently proposed a mechanism for cross linking of P3HB and PLA based on radical formation of DCP at 175 C in melt blending (Fig. 9.14C). However, the thermal characterization of the final polymer composite indicated that cross linking has a negative effect on thermal stability of P3HB due to lowered Tg, Tm, crystallization (%), and spherulite sizes in the final composite (Wei and McDonald, 2015). This result was in contrast with those thermal properties reported by Kolahchi and Kontopoulou (2015) for P3HB/triallyl trimesate composite cross linked by DCP. They observed a significant increase in the Tc and thermal stability of the final composite due to decrease of crystal spherulite sizes and increase in gel fraction and viscosity. It is worth mentioning that other organic-based compounds have been blended into PHAs include dendrimers (Xu et al., 2006), epoxy-functionalized compounds (Duangphet et al., 2014), lignin (Kovalcik et al., 2015; Luo et al., 2016), hydrogen bonding (H-bond) monomer
bisphenol A (Fei et al., 2004b), PEG (Cheng et al., 2006), tannic acid (Xiang et al., 2015), tea polyphenols (Xiang et al., 2015) and monoterpene derivatives, such as linalool (Xiang et al., 2013). These compounds can have either a positive or negative effect on thermal and mechanical properties of PHAs due to their intrinsic properties. The properties of the final PHA composite materials have shown to be highly dependent on how these compounds are cross linked. Incorporation of inorganic or cellulosic-based particles into PHAs through melt coating has also been widely reported by researchers. Liu et al. (2002) used boron nitride, talc, terbium oxide (Tb2O3), and lanthanum oxide (La2O3) as nucleating agents to study the thermal properties of PHBV. They found that during the crystallization, nucleating agents increase the Tc of PHBV compared with that for neat polymer. However, it has little effect on Tm of PHBV after addition of nucleating agents. Calcium carbonate has also been considered as a low-cost filler for PHAs to study the thermal properties (Ding et al., 2011). It was noted that the thermal properties of the final composites can be tailored by calcium carbonate depending on the particle size and the purity. More recent studies on the effect of nanoclays and silver nanoparticles on thermal properties of PHAs have shown that they are not only able to improve the crystallization rate of the polymers but they also increase

Thermal properties of aliphatic polyesters

179

Figure 9.14 (A) Tm and Tc data obtained for PHBV and the cross linked polymer. [The thermal transitions of the samples were determined on a Perkin
Elmer DSC-7c differential scanning calorimeter. The samples were first heated from 20 C to 190 C at a rate of 10 C/min (first heating scan) and, after keeping at 190 C for 1 min, were then cooled to 230 C at a rate of 10 C/min (cooling scan). They were then reheated to 190 C at a rate of 10 C/min (second heating scan). The crystallization temperatures (Tc) were derived from the cooling scans. The melting temperatures (Tm) and heats of fusion were derived from the second heating scans in order to eliminate the effects of different heat histories on the samples. Figure 5 in original paper.] (B) TGA curves of PHBV, PBS, the PHBV/PBS/DCP (80:20:0.5) blend, and gel of the PHBV/PBS/DCP (80:20:0.5) blend. [TGA (Perkin
Elmer, Inc., United States) was used to analyze the gel composition of the PHBV/PBS/DCP blends. PHBV, PBS, PHBV/PBS blends and the extracted residues were heated from room temperature to 700 C in nitrogen atmosphere (40 mL/min) at a heating rate of 20 C/min. Figure 3 in original paper.] (C) The proposed mechanism for cross linking of P3HB and PCL by DCP (a) thermal decomposition of DCP into radicals when exposed to heat (b) reaction of DCP radicals with P3HB and (c) reaction of DCP radicals with PLA (Scheme 1 in original paper). (Continued)

180

Thermal Analysis of Textiles and Fibers

L

the Tg, thermal stability, and storage modulus of the composites depending on the nanoclay content (Naguib et al., 2013; Ciou et al., 2014; Iggui et al., 2015; CastroMayorga et al., 2016). Cellulose nanowhiskers (Ten et al., 2010, 2013), nanofibrillated cellulose (Srithep et al., 2013), and α-cellulose fibers (Wei et al., 2015) have also been introduced into PHAs by solution casting or melt compounding to produce green biocomposites. Ten et al. (2010, 2013) reported that the Tg of PHBV does not change by the addition of various contents of cellulose nanowhiskers while Tc can be decreased in the produced composite due to its nucleation effect. Srithep et al. (2013) observed similar effect but lower thermal stability of the PHBV matrix due to addition of nanofibrillated cellulose. A recent study on the incorporation of α-cellulose fibers and DCP into PHA polymers has shown that cellulosic-based additives and DCP produced cross links between P3HB and PHBV under melt blending. As a result, C
C bonds are formed among P3HB, PHBV, DCP, and cellulose fibers which enhance the thermal stability of the resultant grafted biocomposites (Fig. 9.15) (Wei et al., 2015). Finally, blending of PHAs with other polymers is another method to create new polymer composites with thermal, physical, and mechanical properties, superior than the original PHAs. These polymers included polyvinyl acetate (El-Hadi et al., 2002; Wang et al., 2005; Yang and Hu, 2008), polyethylene vinyl acetate (Guo et al., 2014), and polypropylene carbonate (Li et al., 2004, 2015). On the other hand, researchers performed blending of different individual polymers in PHAs family due to their high compatibility in the mixing process. These blends were also conducted to evaluate either melting or solvent casting affects different properties of the final composites (Dong et al., 2013; Zhao et al., 2003; Kabe et al., 2012). Polyvinyl acetate has been shown as an effective polymer in blending with PHAs to decrease its brittleness (El-Hadi et al., 2002; Wang et al., 2005). Despite improvement in the flexibility of PHAs in blend with polyethylene vinyl acetate, it reduces the Tg and crystallinity (El-Hadi et al., 2002). This result was in contrast with the results obtained by Wang et al. (2005). They found that poly (vinyl acetate) increases the Tg of P3HB/poly(propylene carbonate) blend. They concluded two reasons for this improvement: (1) the ability of poly(vinyl acetate) to work as a compatibilizer between each polymer phase making their chain motions more restrictive; (2) improving the miscibility of P3HB and poly(propylene carbonate) in the blend (Wang et al., 2005). It should be noted that P3HB and poly(propylene carbonate) are immiscible for blends with P3HB content .40% as reported by Yang and Hu (2008). Such improvement in the thermal Source: (A) From Fei, B., Chen, C., Chen, S., Peng, S., Zhuang, Y., An, Y., et al., 2004a. Polym. Int. 53, 937. Reprinted with permission from John Wiley and Sons; (B) From Ma, P., Hristova-Bogaerds, D.G., Lemstra, P.J., Zhang, Y., Wang, S., 2012. Macromol. Mater. Eng. 297, 402. Reprinted with permission from John Wiley and Sons; and (C) From Wei, L., McDonald, A.G., 2015. J. Appl. Polym. Sci. 132, 41724. Reprinted with permission from John Wiley and Sons.

Thermal properties of aliphatic polyesters

181

Figure 9.15 The proposed mechanism for (A) cross linking among P3HB, PHBV, DCP, and cellulose fibers, (B) the chemical structures of grafted αCell-g-P3HB biocomposites, (C) αCell-g-PHBV biocomposites (Figure 1 in original paper). Source: Reported by Wei, L., Stark, N.M., McDonald, A.G., 2015. Green Chem. 17, 4800. Reprinted with permission from the Royal Society of Chemistry.

stability of PHA polymer was also reported by Guo et al. (2014) and was found to be highly dependent on ethylene vinyl acetate content (Fig. 9.16A). It is interesting to note that PHAs acted as nucleating agents in PHA/PLA composites, resulting in formation of PLA lattices and improvement in crystallization rate (Li et al., 2004). The changes in crystallization rate and mechanical properties of these composites are more pronounced when DCP was used as a cross linking agent (Li et al., 2015). As mentioned before, different PHA polymers can be mixed when produced solvent casting or melt blending. In this regard, Zhao et al. (2003) have found that poly(hydroxybutyrate) and poly(hydroxybutyrate-co-hydroxyhexanoate) are miscible due to appearance of only one Tg at 25.15 C while Tc was 20 C (Fig. 9.16B). P3HB with β form crystals can also be mixed with those having lamellar type ones by solvent casting and subsequent cold drawing. This process was found useful for improvement of Tg and Tc of P3HB (Fig. 9.16C). A recent study showed that P3HB can also be mixed with PHO causing an increase in the Tg, Tm, and Tc of the blends in comparison to the neat PHO (Kabe et al., 2012).

182

Thermal Analysis of Textiles and Fibers

Figure 9.16 (A) TGA of P(3HB-co-4HB)/EVA composites (Guo et al., 2014). [TGA Q50 TA. Heating rate was 20 C/min, scanning temperature range was from room temperature to 700 C, and flow atmosphere was under the high purity nitrogen at 50 mL/min. Figure 4(a) in original paper.] (B) DSC curves of polyhydroxybutyrate and poly(hydroxybutyrate-co-hydroxyhexanoate) in comparison with individual polymer. [DSC was performed on 2910 Modulated DSC (TA, United States). Samples (2 mg) were first heated from 25 C to 200 C at a heating rate of 10 C/min (Run 1). After being kept at 200 C for 2 min, the samples were rapidly quenched to 260 C. They were again heated at a rate of 10 C/min to 200 C (Run 2). DSC curves were recorded during run 2. Figure 2 in original paper.] (C) DSC heating curves for P(3HB), 5/95 blend, 10/90 blend, and UHMW-P(3HB) films. [The melting temperature (Tm) and Tg of the blend films were measured by DSC (Perkin
Elmer, DSC8500). The temperature was calibrated using indium. The first heating cycle was carried out at a rate of 20 C/min from 250 C to 200 C, maintained at 200 C for 1 min, and the cooled to 250 C at a rate of 2200 C/min. Tg and Tm were determined from DSC curves obtained for the second heating cycle at a rate of 20 C/min. Figure 1 in original paper.] Source: (A) From Guo, J., Liu, M., Liu, Y., Hu, C., Xia, Y., Zhang, H., and Gong, Y. (2014) J. Appl. Polym. Sci. 131, 41206. Reprinted with permission from John Wiley and Sons (B) From Zhao, K., Deng, Y., Chen, J.C., Chen, G.-Q., 2003. Biomaterials 24, 1041. Reprinted with permission from Elsevier, and (C) From Kabe, T., Tsuge, T., Kasuya, K., Takemura, A., Hikimai, T., Takata, M., et al., 2012. Macromolecules 45, 1858. Reprinted with permission from the American Chemical Society.

Thermal properties of aliphatic polyesters

9.6

183

Conclusion

Biodegradable aliphatic polyesters are of one of the most common family of polymers used in tissue engineering, medical science, and drug delivery. Studies on thermal properties of this family have mostly focused on Tg, Tm, degree of crystallinity, and decomposition from ambient temperature to high temperatures by DSC and TGA. Most thermal degradation starts in the amorphous regions of the aliphatic polyesters before being expanded to the crystalline regions. Using calorimetric techniques to study the thermal properties of aliphatic polyesters opens up several perspectives: 1. It has been seen that there is direct relationship between thermal degradation of aliphatic polyesters and biodegradability. Controlling thermal properties of polyesters may not be simple due to the fact that several parameters can explain the changes in these properties including the polyester type, synthesis method, the structure of main chain and side groups, branching degree, degree of crystallinity, and the type and amount of additives used in blending. 2. New synthesis techniques, copolymers, and additives are needed to improve the thermal properties and biodegradability of aliphatic polyesters for tissue engineering application. 3. Although additives in blending with aliphatic polyesters can improve thermal parameters, the crystallization rate, and thermal stability of polyesters, they may have a negative impact in the environment; therefore, their biocompatibility must be considered.

References Albuerne, J., Marquez, L., Mu¨ller, A.J., Raquez, J.M., Dege´e, P., Dubois, P., et al., 2003. Macromolecules 36, 1633. Ali, S.A.M., Doherty, P.J., Williams, D.F., 1994. J. Appl. Polym. Sci. 51, 1389. Aoyagi, Y., Yamashita, K., Doi, Y., 2002. Polym. Degrad. Stab. 76, 53. Arican, M.O., Mert, O., 2015. RSC Adv. 5, 71519. Arnal, M.L., Lopez-Carrasquero, F., Laredo, E., Mu¨ller, A.J., 2004. Eur. Polym. J. 40, 1461. Arslan, H., Ye¸silyurt, N., Hazer, B., 2007. J. Appl. Polym. Sci. 106, 1742. Ashammakhi, N.A., 1996. J. Biomed. Mater. Res. 33, 297. Baez, J.E., Marcos-Fernandez, A., 2015. Int. J. Polym. Anal. Charact. 20, 637. Balsamo, V., Newman, D., Gouveia, L., Herrera, L., Grimau, M., Laredo, E., 2006. Polymer 47, 5810. Barakat, I., Dubois, Ph, Grandfils, Ch, Je´roˆ me, R., 2001. J. Polym. Sci., A Polym. Chem 39, 294. Becker, J., Lu, L., Runge, M.B., Zeng, H., Yaszemki, M.J., Dadsetan, M., 2015. J. Biomed. Mater. Res., A 103, 2549. Bergsma, J.E., de Bruijn, W.C., Rozema, F.R., Bos, R.R., Boering, G., 1995. Biomaterials 16, 25. Bogdanov, B., Vidts, A., Van den Bulcke, A., Verbeeck, R., Schacht, E., 1998. Polymer 39, 1631. Braun, E., Levin, B.C., 1986. Fire Mater. 10, 107.

184

Thermal Analysis of Textiles and Fibers

Brito, Y., Sabino, M.A., Ronca, G., Albuerne, J., Mu¨ller, A.J., 2006. Rev. Latin Am. Metal. Mater. 26, 61. Brode, G.L., Koleske, J.V., 1972. J. Macromol. Sci.
Chem., A 6, 1109. Castro-Mayorga, J.L., Fabra, M.J., Lagaron, J.M., 2016. Innov. Food Sci. Emerg. Technol. 33, 524. Celli, A., Scandola, M., 1992. Polymer 33, 2699. Cesur, S., Alp, B., Ku¨c¸u¨kgo¨ksel, Y., Kahraman, T., Balko¨se, D., 2015. J. Vinyl Addit. Technol. 21, 174. Chandra, R., Rustgi, R., 1998. Prog. Polym. Sci. 23, 1273. Chen, D.R., Bei, J.Z., Wang, S.G., 2000. Polym. Degrad. Stab. 67, 455. Cheng, G., Wang, T., Zhao, Q., Ma, X., Zhang, L., 2006. J. Appl. Polym. Sci. 100, 1471. Chin, S.J., Vempati, S., Dawson, P., Knite, M., Linarts, A., Ozols, K., et al., 2015. Polymer 58, 209. Ciou, C.Y., Li, S.-Y., Wu, T.M., 2014. Eur. Polym. J. 59, 136. Cipitria, A., Skelton, A., Dargaville, T.R., Dalton, P.D., Hutmacher, D.W., 2011. J. Mater. Chem. 21, 9419. Coleman, N.J., Craig, D.Q.M., 1996. Int. J. Pharm. 135, 13. Coombes, A.G.A., Rizzi, S.C., Williamson, M., Barralet, J.E., Downes, S., Wallace, W.A., 2004. Biomaterials 25, 315. Dalkoni-Veress, K., Forrest, J.A., Murray, C., Gigault, C., Dutcher, J.R., 2001. Phys. Rev. E 63, 1801. Darwis, D., Mitomo, H., Enjoji, T., Yoshii, F., Makuuchi, K., 1998. J. Appl. Polym. Sci. 69, 581. Dash, T.K., Konkimalla, V.B., 2012. J. Controlled Release 158, 15. da Silva, L.G., Leyva, M.E., Barrak, E´.R., Barca, L.F., Sachs, D., de Queiroz, A.A.A., 2011. Synthesis of poly(ε-caprolactone) by iodine: an interesting route to synthesize bioreabsorbable polymers via green chemistry. 11 Congresso Brasileiro de Polimeros 4129
4134. de Koning, G.J.M., Lemstra, P.J., 1993. Polymer 34, 4089. de Koning, G.J.M., Lemstra, P.J., 1994. Polymer 35, 4598. Demirdo¨gen, B., Plazas Bonilla, C.E., Trujillo, S., Perilla, J.E., Elcin, A.E., Elcin, Y.M., et al., 2014. J. Biomed. Mater. Res., A 102, 3229. Dhawan, J.C., Legendre, R.C., Bencsath, A.F., Davis, R.M., 1991. J. Supercrit Fluids 4, 160. Di Lorenzo, M.L., 2005. Eur. Polym. J. 41, 569. Ding, C., Wang, Y., Zhang, S., 2007. Eur. Polym. J. 43, 4244. Ding, Y., He, J., Yang, Y., Cui, S., Xu, K., 2011. Polym. Compos. 32, 1134. Doi, Y., Segawa, A., Kunioka, M., 1990. Int. J. Biol. Macromol. 12, 106. Doi, Y., Kitamura, S., Abe, H., 1995. Macromolecules 28, 4822. Dong, W., Ma, P., Wang, S., Chen, M., Cai, X., Zhang, Y., 2013. Polym. Degrad. Stab. 98, 1549. Doppalapudi, S., Jain, A., Khan, W., Domb, A.J., 2014. Polym. Adv. Technol. 25, 427. Dorgan, J.R., Janzen, J., Clayton, M.P., 2005. J. Rheol. 49, 607. Drumright, R.E., Gruber, P.R., Henton, D.E., 2000. Adv. Mater. 12, 1841. Duangphet, S., Szegda, D., Song, J., Tarverdi, K., 2014. J. Polym. Environ. 22, 1. Dziadek, M., Menaszek, E., Zagrajczuk, B., Pawlik, J., Cholewa-Kowalska, K., 2015a. Mater. Sci. Eng., C 56, 9. Dziadek, M., Pawlik, J., Menaszek, E., Stodolak-Zych, E., Cholewa-Kowalska, K., 2015b. J. Biomed. Mater. Res., B Appl. Biomater. 103, 1580. Ebrahimi, I., Kiumarsi, A., Parvinzadeh Gashti, M., Rashidian, R., Hossein Norouzi, M., 2011. Eur. Phys. J. Appl. Phys. 56, 10801.

Thermal properties of aliphatic polyesters

185

El-Hadi, A., Schnabel, R., Straube, E., Mu¨ller, G., Henning, S., 2002. Polym. Test. 21, 665. Epple, M., Herzberg, O., 1998. J. Biomed. Mater. Res. 43, 83. Erythropel, H.C., Shipley, S., Bo¨rmann, A., Nicell, J.A., Mari´c, M., Leask, R.L., 2016. Polymer 89, 18. Fei, B., Chen, C., Chen, S., Peng, S., Zhuang, Y., An, Y., et al., 2004a. Polym. Int. 53, 937. Fei, B., Chen, C., Wu, H., Peng, S., Wang, X., Dong, L., et al., 2004b. Polymer 45, 6275. Fischer, E.W., Sterzel, H.J., Wegner, G., 1973. Kolloid-Zeitschrift und Zeitschrift fur Polymere 251, 980. Forrest, J.A., Dalnoki-Veress, K., Dutcher, J.R., 1997. Phys. Rev. E 56, 5705. Gan, Z., Liang, Q., Zhang, J., Jing, X., 1997. Polym. Degrad. Stab. 56, 209. Garlotta, D., 2001. J. Polym. Environ. 9, 63. Goto, M., Sasaki, M., Hirose, T., 2006. J. Mater. Sci. 41, 1509. Groot, W., van Krieken, J., Sliekersl, O., de Vos, S., 2010. In: Auras, R., et al., (Eds.), Poly (Lactic Acid): Synthesis, Structures, Properties, Processing, and Applications. John Wiley & Sons, Inc. Guo, J., Liu, M., Liu, Y., Hu, C., Xia, Y., Zhang, H., et al., 2014. J. Appl. Polym. Sci. 131, 41206. Hajiraissi, R., Parvinzadeh, M., 2011. Appl. Surf. Sci. 257, 8443. Hatakeyama, H., Yoshida, T., Hatakeyama, T., 2000. J. Therm. Anal. Calorim. 59, 157. Hazer, B., 2015. Polym. Degrad. Stab. 119, 159. Hazer, B., Steinbu¨chel, A., 2007. Appl. Microbiol. Biotechnol. 74, 1. Herrera, N., Salaberria, A.M., Mathew, A.P., Oksman, K., 2016. Compos., A 83, 89. Herweh, J.E., Work, J.L., Whitmore, W.Y., 1977. The dependence of the dynamic mechanical properties of poly[(1,4-cyclohexane)bismethylene isophthalate]-b-polycaprolactone copolymers on segment size and crystallinity. In: Klempner, D., Frisch, K.C. (Eds.), Polymer Alloys. Polymer Science and Technology, vol 10. Springer, Boston, MA. Hoogsteen, W., Postema, A.R., Pennings, A.J., Brinke, G.T., Zugenmaier, P., 1990. Macromolecules 23, 634. Hoque, M.E., San, W.Y., Wei, F., Li, S., Huang, M.-H., Vert, M., et al., 2009. Tissue Eng., A 15, 3013. Huang, Z.-M., Zhang, Y.-Z., Kotaki, M., Ramakrishna, S., 2003. Compos. Sci. Technol. 63, 2223. Hyon, S.H., Jamshidi, K., Ikada, Y., 1984. Spinning of poly-L-lactide and hydrolysis of the fiber in vitro. In: Shalaby, S.W., Hoffman, A.S., Ratner, B.D., Horbett, T.A. (Eds.), Polymers as Biomaterials. Plenum Press, New York, pp. 51
66. Iggui, K., Le Moigne, N., Kaci, M., Cambe, S., Degorce-Dumas, J.-R., Bergeret, A., 2015. Polym. Degrad. Stab. 119, 77. Ikada, Y., Tsuji, H., 2000. Macromol. Rapid Commun. 21, 117. Jayakumar, R., Tamura, H., 2008. Int. J. Biol. Macromol. 43, 32. Jerome, C., Lacomte, P., 2008. Adv. Drug Deliv. Rev. 60, 1056. Jia, Y., Shen, X., Gu, X., Dong, J., Mu, C., Zhang, Y., 2008. Polym. Adv. Technol. 19, 159. Jing, F., Smith, M.R., Baker, G.L., 2007. Macromolecules 40, 9304. Joshi, P., Madras, G., 2008. Polym. Degrad. Stab. 93, 1901. Kabe, T., Tsuge, T., Kasuya, K., Takemura, A., Hikimai, T., Takata, M., et al., 2012. Macromolecules 45, 1858. Kawamoto, N., Sakai, A., Horikoshi, T., Urushihara, T., Tobita, E., 2007. J. Appl. Polym. Sci. 103, 198. Ke, Y., Zhang, X.Y., Ramakrishna, S., He, L.M., Wu, G., 2017. Mater. Sci. Eng., C 70, 1107. Kesenci, K., Fambri, L., Migliaresi, C., Piskin, E., 2000. J. Biomater. Sci., Polym. Ed. 11, 617.

186

Thermal Analysis of Textiles and Fibers

Kesenci, K., Motta, A., Fambri, L., Migliaresi, C., 2001. J. Biomater. Sci., Polym. Ed. 12, 337. Khang, G., Choee, J.H., Rhee, J.M., Lee, H.B., 2002. J. Appl. Polym. Sci. 85, 1253. Kim, H.W., Chung, M.G., Kim, Y.B., Rhee, Y.H., 2008. Int. J. Biol. Macromol. 43, 307. Kimura, Y., Matsuzaki, Y., Yamane, H., Kitao, T., 1989. Polymer 30, 1342. Kobayashi, S., Uyama, H., Takamoto, T., 2000. Biomacromolecules 1, 3. Kolahchi, A.R., Kontopoulou, M., 2015. Polym. Degrad. Stab. 121, 222. Koronis, G., Silva, A., Fontul, M., 2013. Compos., B 44, 120. Kovalcik, A., Machovsky, M., Kozakova, Z., Koller, M., 2015. React. Funct. Polym. 94, 25. Kulizinski, Z., Piorkowska, E., 2005. Polymer 46, 10290. Kunioka, M., Kawaguchi, Y., Doi, Y., 1989. Appl. Microbiol. Biotechnol. 30, 569. Kuo, Y.C., Leou, S.N., 2006. Biotechol. Prog. 22, 1664. Kurcok, P., Penczek, J., Franek, J., Jedlinski, Z., 1992. Macromolecules 25, 2285. Kweon, D.-K., Cha, D.-S., Park, H.-J., Lim, S.-T., 2000. J. Appl. Polym. Sci. 78, 986. Kwon, I.K., Park, K.D., Choi, S.W., Lee, S.-H., Lee, E.B., Na, J.S., et al., 2001. J. Biomater. Sci., Polym. Ed. 12, 1147. Labet, M., Thielemans, W., 2009. Chem. Soc. Rev. 38, 3484. Lee, J.H., Park, T.G., Park, H.S., Lee, D.S., Lee, Y.K., Yoon, S.C., et al., 2003a. Biomaterials 24, 2772. Lee, S.H., Kim, B.-S., Kim, S.H., Choi, S.W., Jeong, S.I., Kwon, I.K., et al., 2003b. J. Biomed. Mater. Res., A 66A, 29. Lee, K.M., Knight, P.T., Chung, T., Mather, P.T., 2008. Macromolecules 41, 4730. Li, J., Lai, M.F., Liu, J.J., 2004. J. Appl. Polym. Sci. 92, 2514. Li, H., Huneault, M.A., 2007. Polymer 48, 6855. Li, H., Lu, X., Yang, H., Hu, J., 2015. J. Therm. Anal. Calorim. 122, 817. Lim, L.T., Auras, R., Rubino, M., 2008. Prog. Polym. Sci. 33, 820. Lin, W.J., 1999. J. Biomed. Mater. Res., A 47, 420. Linhart, W., Peters, F., Lehmann, W., Schwarz, K., Schilling, A.F., Amling, M., et al., 2001. J. Biomed. Mater. Res. 54, 162. Liu, W.J., Yang, H.L., Wang, Z., Dong, L.S., Liu, J.J., 2002. J. Appl. Polym. Sci. 86, 2145. Lovinger, A.J., Han, B.J., Padden Jr., F.J., Mirau, P.A., 1993. J. Polym. Sci., B: Polym. Phys. 31, 115. Luetzow, K., Klein, F., Weigel, T., Apostel, R., Weiss, A., Lendlein, A., 2007. J. Biomech. 40, S80. Lunt, J., 1998. Polym. Degrad. Stab. 59, 145. Luo, S., Cao, J., McDonald, A.G., 2016. ACS Sustain. Chem. Eng. 4, 3465. Luzier, W.D., 1992. Proc. Natl. Acad. Sci. U.S.A. 89, 839. Ma, P., Hristova-Bogaerds, D.G., Lemstra, P.J., Zhang, Y., Wang, S., 2012. Macromol. Mater. Eng. 297, 402. Macrae, R.M., Wilkinson, J.F., 1958. J. Gen. Microbiol. 19, 210. Maglio, G., Migliozzi, A., Palumbo, R., 2003. Polymer 44, 369. Maharana, T., Mohanty, B., Negi, Y.S., 2009. Prog. Polym. Sci. 34, 99. Mark, J.E. (Ed.), 1999. Polymer Data Handbook. Oxford University Press, New York. Marquez, Y., Franco, M., Turon, P., Puiggalı´, J., 2014. Thermochim. Acta 585, 71. Marquez, Y., Franco, M., Martı´nez, J.C., Estrany, F., Turon, P., Puiggalı´, J., 2015. Eur. Polym. J. 73, 222. Mathias, L.J., Kusefoglu, S.A., Kress, A.O., 1989. J. Appl. Polym. Sci. 38, 1037. Matzinos, P., Tserki, V., Gianikouris, C., Pavlidou, E., Panayiotou, C., 2002. Eur. Polym. J. 38, 1713. Mazzullo, S., Paganetto, G., Celli, A., 1992. Prog. Colloid Polym. Sci. 87, 32.

Thermal properties of aliphatic polyesters

187

Megna, I.S., Koroscil, A., 1968. J. Polym. Sci., C: Polym. Lett. 6, 653. Menczel, J.D., Collins, G.L., Saw, S.K., 1997. J. Therm. Anal. 49, 201. Migliaresi, C., Cohn, D., De Lollis, A., Fambri, L., 1991a. J. Appl. Polym. Sci. 43, 83. Migliaresi, C., De Lollis, A., Fambri, L., Cohn, D., 1991b. Clin. Mater. 8, 111. Min, C., Cui, W., Bei, J., Wang, S., 2005. Polym. Adv. Technol. 16, 608. Mu¨ller, A.J., Arnal, M.L., 2005. Prog. Polym. Sci. 30, 559. Naguib, H.F., Abdel Aziz, M.S., Saad, G.R., 2013. J. Ind. Eng. Chem. 19, 56. Najemi, L., Jeanmaire, T., Zerroukhi, A., Raihane, M., 2010. Starch 62, 147. Nejad, E.H., Paoniasari, A., Koning, C.E., Duchateau, R., 2012. Polym. Chem. 3, 1308. Ouhadi, T., Stevens, C., Teyssie´, P., 1976. J. Appl. Polym. Sci. 20, 2963. Ozawa, T., 2000. Thermochim. Acta 355, 35. Park, K.E., Kang, H.K., Lee, S.J., Min, B.M., Park, W.H., 2006. Biomacromolecules 7, 635. Parvinzadeh, M., Hajiraissi, R., 2007. AIP Conf. Proc. 929, 216. Parvinzadeh, M., Hajiraissi, R., 2008. Tensile Surfactants Deterg. 45, 254. Parvinzadeh Gashti, M., Willoughby, J., Agrawal, P., 2011. Surface and bulk modification of synthetic textiles to improve dyeability. In: Hauser, P.J. (Ed), Textile Dyeing. Available from: ,http://www.intechopen.com/books/textile-dyeing. Parvinzadeh Gashti, M., Moradian, S., 2012. J. Appl. Polym. Sci. 125, 4109. Parvinzadeh Gashti, M., Parvinzadeh Gashti, M., 2012. J. Dispersion Sci. Technol. 34, 853. Parvinzadeh Gashti, M., Yousefpour Navid, M., Rahimi, M.H., 2012. J. Appl. Polym. Sci. 125, 1430. Parvinzadeh Gashti, M., Moradian, S., Rashidi, A., Yazdanshenas, M.E., 2013a. Fibers Polym. 14, 743. Parvinzadeh Gashti, M., Yousefpour Navid, M., Rahimi, M.H., 2013b. Pigm. Resin Technol. 42, 34. Parvinzadeh Gashti, M., Hajiraissi, R., Parvinzadeh, M., 2013c. Fibers Polym. 14, 1324. Parvinzadeh Gashti, M., Hegemann, D., Stir, M., Hulliger, M., 2014. Plasma Process. Polym. 11, 37. Parvinzadeh Gashti, M., Ebrahimi, I., Pousti, M., 2015. Curr. Appl. Phys. 15, 1075. Parvinzadeh Gashti, M., Stir, M., Hulliger, J., 2016. New J. Chem. 40, 5495. Paul, M.A., Alexandre, M., Dege´e, P., Henrist, C., Rulmont, A., Dubois, P., 2003. Polymer 44, 443. Paul, S., Zhu, Y., Romain, C., Brooks, R., Saini, P.K., Williams, C.K., 2015. Chem. Commun. 51, 6459. Perego, G., Cella, G.D., Bastioli, C., 1996. J. Appl. Sci. 59, 37. Persenaire, O., Alexandre, M., Dege´e, P., Dubois, P., 2001. Biomacromolecules 2, 288. Petersson, L., Kvien, I., Oksman, K., 2007. Compos. Sci. Technol. 67, 2535. Pistner, H., Bendix, D.R., Mu¨hling, J., Reuther, J.F., 1993. Biomaterials 14, 291. Pitt, C.G., 1990. Polycaprolactone and its copolymers. In: Chasin, M., Langer, R. (Eds.), Biodegradable Polymers as Drug Delivery Systems. Marcel Dekker, New York, p. 71. Pitt, C.G., Chasalow, F.I., Hibionada, Y.M., Klimas, D.M., Schindler, A., 1981. J. Appl. Polym. Sci. 26, 3779. Plazas Bonilla, C.E., Trujillo, S., Demirdo¨gen, B., Perilla, J.E., Elcin, Y.M., Gomez Ribelles, J.L., 2014. Mater. Sci. Eng., C 40, 418. Rahimi, M.H., Parvinzadeh, M., Yousefpour Navid, M., Ahmadi, S., 2011. J. Surfactants Deterg. 14, 595. Ray, S.S., Yamada, K., Okamoto, M., Ueda, K., 2003. Polymer 44, 857.

188

Thermal Analysis of Textiles and Fibers

Rojas de Gascue, B., Mendez, B., Manosalva, J.L., Ruiz Santa Quiteria, V., Mu¨ller, A.J., 2002. Polymer 43, 2151. Roushandeh, J.M., Nabi Sarbolouki, M., 2001. Iran. Polym. J. 10, 53. Ryner, M., Albertsson, A.C., 2002. Biomacromolecules 3, 601. Sabino, M.A., 2007. Polym. Degrad. Stab. 92, 986. Sabino, M.A., Feijoo, J.L., Mu¨ller, A.J., 2001. Polym. Degrad. Stab. 73, 541. Saito, Y., Doi, Y., 1994. Int. J. Biol. Macromol. 16, 99. Savelyeva, X., Li, L., Mari´c, M., 2015. J. Polym. Sci., A: Polym. Chem. 53, 59. Schlesinger, E., Ciaccio, N., Desai, T.A., 2015. Mater. Sci. Eng., C 57, 232. Schmidt, C., Behl, M., Lendlein, A., Beuermann, S., 2014. RSC Adv. 4, 35099. Schwarz, K., Epple, M., 1999. Macromol. Chem. Phys. 200, 2221. Seefried Jr, C.G., Koleske, J.V., Critchfield, F.E., 1975. J. Appl. Polym. Sci. 19, 2493. Seyednejad, H., Ghassemi, A.H., van Nostrum, C.F., Vermonden, T., Hennink, W.E., 2011. J. Controlled Release 152, 168. Shakoor, A., Muhammad, R., Thomas, N.L., Silberschmidt, V.V., 2013. J. Phys. Conf. Ser. 451, 1. Sinha, V.R., Bansal, K., Kaushik, R., Kumria, R., Trehan, A., 2004. Int. J. Pharm. 278, 1. Sisson, A.L., Ekinci, D., Lendlein, A., 2013. Polymer 54, 4333. Sivalingam, G., Madras, G., 2003. Polym. Degrad. Stab. 80, 11
16. Sivalingam, G., Karthik, R., Madras, G., 2003a. J. Anal. Appl. Pyrolysis 70, 631. Sivalingam, G., Chattopadhyay, S., Madras, G., 2003b. Polym. Degrad. Stab. 79, 413. Sivalingam, G., Vijayalakshmi, S.P., Madras, G., 2004. Ind. Eng. Chem. Res. 43, 7702. Sodergard, A., Stolt, M., 2002. Prog. Polym. Sci. 27, 1123. Song, P., Jiang, S., Ren, Y., Zhang, X., Qiao, T., Song, X., et al., 2016. J. Colloid Interface Sci. 479, 160. Spearman, S.S., Rivero, I.V., Abidi, N., 2014. J. Appl. Polym. Sci. 131, 40224. Spearman, S.S., Irin, F., Rivero, I.V., Green, M.J., Abidi, N., 2015. Polymer 56, 476. Sravanthi, R., Pramanik, K., Pal, K., 2010. Int. J. Biol. Sci. Eng. 01, 195. Srithep, Y., Ellingham, T., Peng, J., Sabo, R., Clemons, C., Turng, L.-S., et al., 2013. Polym. Degrad. Stab. 98, 1439. Stridsberg, K., Albertsson, A.C., 2000. J. Polym. Sci., A: Polym. Chem. 38, 1774. Sudesh, K., Abe, H., Doi, Y., 2000. Prog. Polym. Sci. 25, 1503. Sun, H., Mei, L., Song, C., Cui, X., Wang, P., 2006. Biomaterials 27, 1735. Tan, H.W., Abdul Aziz, A.R., Aroua, M.K., 2013. Renew. Sustain. Energy Rev. 27, 118. Ten, E., Turtle, J., Bahr, D., Jiang, L., Wolcott, M., 2010. Polymer 51, 2652. Ten, E., Jiang, L., Wolcott, M.P., 2013. Carbohydr. Polym. 92, 206. Tokiwa, Y., Calabia, B.P., 2004. Biotechnol. Lett. 26, 1181. Tokiwa, Y., Calabia, B.P., 2007. J. Polym. Environ. 15, 259. Trujillo, M., Arnal, M.L., Mu¨ller, A.J., Mujica, M.A., Urbina de Navarro, C., Ruelle, B., et al., 2012. Polymer 53, 832. Tsuji, H., Tashiro, K., Bouapao, L., Narita, J., 2008. Macrom. Mater. Eng. 293, 947. Ugartemendia, J.M., Mun˜oz, M.E., Sarasua, J.R., Santamaria, A., 2014. Rheol. Acta 53, 857. Van Bogart, J.C.W., Gibson, P.E., Cooper, S.L., 1983. J. Polym. Sci., B: Polym. Phys 21, 65. Van de Velde, K., Kiekens, P., 2002. Polym. Test. 21, 433. Vasanthakumari, R., Pennings, A.J., 1983. Polymer 24, 175. Vert, M., 2005. Biomacromolecules 6, 538. Vogel, R., T¨andler, B., Voigt, D., Jehnichen, D., H¨außler, L., Peitzsch, L., et al., 2007. Macromol. Biosci. 7, 820. Wallen, L.L., Rohwedder, W.K., 1974. Environ. Sci. Technol. 8, 576.

Thermal properties of aliphatic polyesters

189

Wang, Z.G., Hsiao, B.S., Zong, X.-H., Yeh, F., Zhou, J.J., Dromier, E., et al., 2000. Polymer 41, 621. Wang, X., Peng, S., Dong, L., 2005. Colloid Polym. Sci. 284, 167. Wang, B., Zheng, X., Chang, M.-W., Ahmad, Z., Li, J.-S., 2016. Colloids Surf., B 145, 757. Wei, L., McDonald, A.G., 2015. J. Appl. Polym. Sci. 132, 41724. Wei, L., Stark, N.M., McDonald, A.G., 2015. Green Chem. 17, 4800. Witzke, D.R., 1997. Introduction to Properties, Engineering, and Prospects of Polylactide Polymer (PhD thesis). Chemical Engineering, Michigan State University. Woodruff, M.A., Hutmacher, D.W., 2010. Progress Polym. Sci. 35, 1217. Wouwe, P.V., Dusselier, M., Vanleeuw, E., Sels, B., 2016. ChemSusChem 9, 907. Wu, C.S., 2004. J. Appl. Polym. Sci. 92, 1749. Wu, T.W., Wu, C.Y., 2006. Polym. Degrad. Stab. 91, 2198. Wu, Q., Wang, L., Fu, Song, X., Yang, Q., Zhang, G., 2014. Polym. Bull. 71, 1. Xiang, H.X., Chen, S.H., Cheng, Y.H., Zhou, Z., Zhu, M.F., 2013. eXPRESS Polym. Lett. 7, 778. Xiang, H., Li, L., Wang, S., Wang, R., Cheng, Y., Zhou, Z., et al., 2015. Polym. Compos. 36, 2303. Xiong, X.Y., Tam, K.C., Gan, L.H., 2006. J. Appl. Polym. Sci. 100, 4163. Xu, S., Luo, R., Wu, L., Xu, K., Chen, G.-Q., 2006. J. Appl. Polym. Sci. 102, 3782. Yang, D.Z., Hu, P., 2008. J. Appl. Polym. Sci. 109, 1635. Yeh, J.-T., Yang, M.-C., Wu, C.-J., Wu, C.-S., 2009. J. Appl. Polym. Sci. 112, 660. Yeo, S.D., Kiran, E., 2005. J. Supercrit. Fluids 34, 287. Yoon, C.S., Ji, D.S., 2005. Fibers Polym. 6, 13. You, Y., Min, B.M., Lee, S.J., Lee, T.S., Park, W.H., 2005. J. Appl. Polym. Sci. 95, 193. Yu, Y.-H., Fan, C.-L., Hsu, Y.-H., Chou, Y.-C., Ueng, S.W.N., Liu, S.-J., 2015. Materials 8, 7714. Zhang, Y., Chen, Z., Dong, Z., Zhao, M., Ning, S., He, P., 2013. Int. J. Polym Mater. 62, 397. Zhao, K., Deng, Y., Chen, J.C., Chen, G.-Q., 2003. Biomaterials 24, 1041. Zhou, S., Xu, J., Yang, H., Deng, X., 2004. Macrom. Mater. Eng. 289, 576. Zia, K.M., Noreen, A., Zuber, M., Tabasum, S., Mujahid, M., 2016. Int. J. Biol. Macrom. 82, 1028.

Further reading Pilon, L., Kelly, C., 2016. J. Appl. Polym. Sci. 133, 42588.

Poly(ethylene naphthalate) [poly (ethylene-2,6-naphthalene dicarboxylate)]

10

Joseph D. Menczel Thermal Measurements LLC, Fort Worth, TX, United States

Abstract Poly(ethylene naphthalate) [poly(ethylene-2,6-naphthalene dicarboxylate), PEN] fibers were introduced recently in areas where high tensile strength and high modulus is necessary. It is more expensive than poly(ethylene terephthalate) (PET), but these fibers have better dimensional stability and other properties. Like for other semicrystalline fibers, the melting point, the heat of fusion, and the tensile modulus of these fibers increases with increasing draw ratio.

Poly(ethylene naphthalate) [poly(ethylene-2,6-naphthalene dicarboxylate), PEN] is a relatively new polyester fiber. This polymer has very similar structure to PET, but the phenylene ring is replaced by a naphthalene ring. It is made by polycondensation of naphthalene dicarboxylic acid (NDA) and ethylene glycol. PEN was very expensive in the 1990s and early 2000s. The situation right now is better because Amoco commercialized NDA, which is the essential raw material of PEN. Later Teijin and Mitsubishi stepped in with the production of the raw material. The major application of PEN is in the area of specialty films (flexible food packaging, electrical, solar cell protection, and pressure sensitive tapes), but it is also used in the fiber areas where high tensile strength and modulus are needed. PEN has a number of advantages over PET: higher temperature resistance, higher temperature of shrinkage, long-term electrical use at high temperatures (up to 155
C), and better resistance to UV as well as hydrolysis in hot humid conditions. The barrier properties (including the oxygen barrier) of PEN are really good, so it is widely used in beverage bottle production where Thermal Analysis of Textiles and Fibers. DOI: https://doi.org/10.1016/B978-0-08-100572-9.00010-0 © 2020 Elsevier Ltd. All rights reserved.

192

Thermal Analysis of Textiles and Fibers

the beverage is easily oxidized (such as beer). This application is one of the major uses of PEN, especially in Europe for bottling beer. The dimensional and thermal stability of PEN is better than that of PET (no shrinkage up to 180
C). Its application is expected to grow fast, and in the near future it can replace classical plastic scintillators, because its scintillation properties are very good (scintillation is a property similar to luminescence). The tenacity of PEN fiber is about 10% higher than that of poly(ethylene terephthalate) (PET), the modulus of PEN fiber is three times as high as that of PET, while its elongation is only half of PET. The major disadvantage of PEN is its higher price over PET. The use of PEN is expected to grow considerably in the film shape in manufacturing of electronic parts such as capacitors, magnetic tapes, and printed circuit boards. While the performance of PEN as a reinforcing fiber in tires is not as superior as of liquid crystalline fibers (e.g., Vectran), it may have some advantages because of its cheaper price. The PEN fiber is made by melt spinning (of course, with postdrawing) similar as PET. The market size of poly(ethylene naphthalate) is quite considerable, by 2020 it is expected to be $1.4 billion (Grand View Research, Inc.). The major users of PEN will be the beverage bottles and the electrical industry. The growth of using PEN in the electrical industry will be propelled by increasing use of PEN in wires, cables, solar cells, and various sensors. The major advantages of PEN over PET are its better thermal stability, higher glass transition temperature, higher dimensional stability (e.g., in tire reinforcement), excellent barrier properties, and better shrinkage properties. Most new applications will be found in the food industry, solar cells, and various sensors. The major disadvantage of PEN is its higher price over PET. The most important companies in the PEN market are Toray, Sumitomo Chemical, DuPont, DuraFiber, and Seiwa. The first patents on melt spinning of PEN belong to Shima et al. (1973), Hamana et al. (1974), and Kumakawa et al. (1975a,b). The most important application of PEN fiber is as reinforcing fiber in rubber tires and manufacturing laminated sailcloth (where the fibers are sandwiched between polymer film sheets to provide increased stability). These applications of PEN are justified by its higher resistance to wear and much higher tenacity than PET fiber. PEN fiber is not cheap and may hinder its application as a tire cord. But its use may be justified by the higher glass transition temperature: during driving the temperature of the tire raises significantly (at about 90 km/hour speed the tire temperature reaches 70
C, which is close to Tg of PET). There are surprisingly few papers on thermal analysis of PEN fibers. Even books such as the Handbook of Thermoplastics (Olabisi and Kolapo, 2016) or High Performance Polymers (Fink, 2008, 2014) contain very limited information on PEN fibers. There are some papers on electrospinning of PEN (e.g., Reneker and Chun, 1996; Reneker et al., 2000; Bognitzki et al., 2001; Kim and Reneker, 1999), but detailed thermal analysis is still missing. Zachmann et al. (1985) studied crystallization kinetics of PEN. He received Tg 5 120
C, and the nonequilibrium melting point of PEN as 280
C. Zachmann found out that the general thermal properties of PEN are very similar to the thermal

Poly(ethylene naphthalate) [poly(ethylene-2,6-naphthalene dicarboxylate)]

193

properties of PET: PEN can also be quenched easily to the fully amorphous state. Now we know that this quenching must take place from temperatures much higher than the equilibrium melting point, because PEN may have a mesophase. In this quasicrystalline state the macromolecular segments were shown not to be fully crystallized in the equatorial direction in the wide angle X-ray scattering (WAXS) experiments, but they were in registry with each other in the meridional direction. In addition, it should be mentioned that PEN has two crystal forms, and both of them are triclinic. The crystal form you get, depends on the crystallization temperature, and it may also depend on the temperature where the cooling starts. If crystallizing at 180
C, the α-form is obtained, while from 240
C, the β-form can be seen. Zachmann et al. (1985) estimated the heat of fusion of 100% crystalline PEN to be ΔHf 5 190 J/g. Cheng and Wunderlich (1988) reported ΔCp 5 80.1 J/(mol
C) for the heat capacity increase in the glass transition, and they estimated the equilibrium
melting point of PEN to be Tm 5 337
C. There are surprisingly few papers on thermal analysis of PEN fibers, perhaps due to the high cost of the basic material. Saw et al. (1997) performed a detailed X-ray and thermal analysis study on asspun and hot drawn PEN fibers. Fig. 10.1 shows the DSC traces of the as-spun PEN fiber (top curve), and two drawn fiber runs (draw ratio, DR 5 5.4 3 ) recorded at 10
C/min heating rate in free-to-shrink and constrained state measurements. The as-spun fiber (i.e., the amorphous polymer) had a glass transition temperature of Tg 5 125
C. The starting temperature of cold crystallization was Tcco 5 175
C, and the melting point was Tm 5 275
C. Crystallinity is increased from 45% to 50% as the draw ratio increased from 2.3 to 5.4. In the constrained state experiments, the melting point increases with draw ratio up to a draw ratio of 4.3 3 and then stays constant at 289
C. The heat of fusion of the drawn fibers is higher when measured in the free-to-shrink state than in the constrained state: for a draw ratio of 5.4, the heats of fusion for the free-to-shrink and the fixed-length measurements were 100.7 and 95.6 J/g. This is very similar to the results obtained by Smook and Pennings (1984) for gel-spun polyethylene fibers, but the difference in the heats of fusion is smaller for PEN than for PE. Smook and Pennings (1984) suggested that the heat of fusion difference for gel-spun polyethylene was due to the orthorhombic-topseudohexagonal transition, and that difference was up to 60 J/g (at the highest draw ratio of DR 5 100 3 , which is a reasonable number for such an effect). In PEN the difference is considerably smaller, and perhaps this can be attributed to the existence of a nematic phase. Indeed, the experiments of Jakeways et al. (1996) and Saw et al. (1997) suggested that some mesophase may exist for PEN: as mentioned earlier, in that state the macromolecular segments are in registry with each other in the meridional direction, but not fully crystallized in the equatorial direction. There is another significant difference between the free-to-shrink and fixedlength DSC measurements of the PEN fibers. When the PEN fibers are melted in the constrained state, and then crystallized and reheated again, (1) the crystallization temperature increases with increasing draw ratio, and (2) the heat of fusion on

194

Thermal Analysis of Textiles and Fibers

As-spun PEN fiber

–0.5

–1.5 Drawn fiber, constrained

–2.5

Endotherm

Heat flow (W/g)

Drawn fiber, free-to-shrink

–3.5 100

150

200

250

300

350

Temperature (ºC) Figure 10.1 DSC curves of as-spun and drawn PEN fibers (DR 5 4.8) recorded in free-to-shrink and constrained measurements. The measurements were performed using a TA Instruments 2100/910 DSC instrument. Heating rate was 10
C/min (Saw et al., 1997). Wellvisible glass transition and cold crystallization can be observed on the DSC curve of the as-spun (amorphous) fiber. The glass transition is blurred on the DSC curve of the drawn fiber, there is no cold crystallization, and the melting point is higher than for the as-spun fiber. DuPont 2200/910 DSC, heating rate 10
C/min. Source: From Saw, C.K., Menczel, J., Choe, E.W., Hughes, O.R., 1997. SPE Antec0 97, Toronto, ON, Canada, p. 1610, Figs. 5, 6, 7. Reprinted with permission of SPE Antec.

the reheating is considerably higher for the constrained samples than for the free-toshrink samples obviously due to the nucleating effect of the remaining orientation in the constrained samples. Fig. 10.2 shows the DMA (E0 vs temperature) curves of the as-spun and drawn PEN fibers. The as-spun fiber has considerably lower modulus, and it shows a significant (almost three orders of magnitude) drop in the modulus between 100
C and 140
C (the glass transition is 125
C). Then at 160
C the modulus suddenly increases (this is the cold crystallization), and finally it decreases slowly again due to melting. The DMA curves of the drawn fibers are completely different: the modulus exhibits smooth, slow decrease with increasing temperature, then there is a more significant drop in the modulus between 140
C and 170
C due to the glass transition. It is noteworthy that the glass transition of the drawn fibers is much

Poly(ethylene naphthalate) [poly(ethylene-2,6-naphthalene dicarboxylate)]

195

Figure 10.2 The tensile storage modulus of as-spun and two drawn PEN fibers with draw ratios of 2.3 and 5.4. For the as-spun fiber the modulus decreases between 120
C and 165
C corresponds to cold crystallization of the amorphous polymer. Tg of the drawn fibers is much higher as can be seen from the slight modulus decrease for the 2.3 3 and 5.4 3 drawn fibers. Also, the modulus decreases at Tg for the drawn fibers is much smaller than for the as-spun fiber. Rheometrics RSA2 DMA, 3
C temperature step method. Source: From Saw, C.K., Menczel, J., Choe, E.W., Hughes, O.R., 1997. SPE Anten 0 97, Toronto, ON, Canada, p. 1610, Figs. 57. Reprinted with permission of SPE Antec.

higher than Tg of the as-spun (amorphous) fiber. This is a very frequent phenomenon: crystallinity raises the glass transition temperature. Fig. 10.3. shows the loss tangent from the DMA measurements. At around 170
C, there is a peak corresponding to the glass transition of PEN, and another peak can be seen at around 60
C: this peak can be assigned to the onset of rotation of the naphthalene rings. The measurements were stopped at 200
C, so it is unknown whether a higher temperature α-relaxation exists in PEN: the knowledge of this transition would be important, because for efficient drawing the draw temperature should be beyond the temperature of the α-relaxation. Kim and Lee (2000) studied the PEN fiber samples before and after electrospinning and came to a conclusion that a significant decrease of the molecular mass occurs during electrospinning. It is not known whether the humidity was controlled in these experiments. If not, a behavior similar to that of PET may take place, that is, chain scission (this is the process essentially used to recycle PET: hydrolytic depolymerization at high temperatures).

196

Thermal Analysis of Textiles and Fibers

0.1

Tan δ

PEN single fiber, DR=4.86x

0.01 –50

0

50

100

150

200

Temperature (ºC)

Figure 10.3 Loss tangent of drawn PEN single fiber as a function of temperature. Rheometrics RSA2 DMA, 3
C temperature step method. Source: From Saw, C.K., Menczel, J., Choe, E.W., Hughes, O.R., 1997. SPE Anten 0 97, Toronto, ON, Canada, p. 1610, Figs. 57. Reprinted with permission of Society of Plastics Engineers.

References Bognitzki, M., Czado, W., Frese, T., Schaper, A., Hellwig, M., Steinhart, M., Greiner, A., Wendorff, J.H., 2001. Adv. Mater. 13 (1), 70. Cheng, S.Z.D., Wunderlich, B., 1988. Macromolecules 21, 789. Fink, J.K., 2014. High Performance Polymers. Elsevier. Hamana, I., Fujiwara, R., Kumagawa, S., 1974. Jpn. Pat. 49,006,771. Kim, J., Lee, D.S., 2000. Polym. J. 32, 616. Kim, J.S., Reneker, D.H., 1999. Polym. Comp. 20, 124. Kumakawa, S., Hamana, I., Fujiwara, Y., 1975a. Jpn. Pat. 50,046,922. Kumakawa, S., Hamana, I., Fujiwara, Y., 1975b. Jpn. Pat. 50,046,923. Olabisi, O., Kolapo, A., 2016. Handbook of Thermoplastics, second ed. CRC Press. Reneker, D.H., Chun, I., 1996. Nanotechnology 7, 216. Reneker, D.H., Yarin, A.L., Fong, H., Koombhonge, S., 2000. J. Appl. Phys. 87, 4531. Saw, C.K., Menczel, J., Choe, E.W., Hughes, O.R., 1997. SPE ANTEC0 97, Toronto, ON, Canada, p. 1610. Shima, T., Yamashiro, S., Yoshimura, M., 1973. Jpn. Pat. 48,010,494. Smook, J., Pennings, J., 1984. Colloid Polym. Sci. 262, 712. Zachmann, H.G., Wiswe, D., Gehrke, R., Rickel, C., 1985. Makromol. Chem., Suppl. 12, 175.

Polyethylene fibers

11

Joseph D. Menczel1 and Tonson Abraham2 1 Thermal Measurements LLC, Fort Worth, TX, United States, 2Retired from ExxonMobil Chemical Company, Avon, OH, United States

Abstract Polyethylene (PE) is the simplest and one of the real, very few homopolymers. Polyethylene fibers, especially the gel-spun fibers, are one of the materials created by mankind having the best mechanical properties. In this chapter the thermal properties obtained by differential scanning calorimetry (DSC) and dynamic mechanical analysis (DMA) are presented. The long-debated glass transition temperature is discussed, and it is shown that modulated temperature thermomechanical analysis (MT-TMA) measurements carried out on Dyneema fibers give 220
C as Tg of PE in agreement with Cp-measurements of the Wunderlich group. The unconstrained and constrained melting experiments of gel-spun PE-fibers are presented. However, no explanation could be given to the highintensity γ-transition in the DMA measurements.

Polyethylene (PE) is the only polymer with known equilibrium melting parameters. Wunderlich et al. prepared equilibrium (i.e., macroscopic extended chain crystals) by high-pressure solution crystallization (Wunderlich and Czornyj, 1977; Prime and Wunderlich, 1969; Prime et al., 1969). From these experiments, we know that the
equilibrium melting point of PE is Tm 5 141:4
C, and the heat of fusion of 100%
crystalline PE is ΔHf 5 293 J=g (4.11 kJ/mol). However, the glass transition of PE is still debated. The Wunderlich school determined 236
C using absolute heat capacity measurements [at the glass transition ΔCp 5 10.3 J/(K mol)] (Gaur and Wunderlich, 1980), but Beatty and Karasz (1979) favor 2128
C by the simple reason that this is the temperature where you can see the most intense peak in the tan δ versus temperature in the dynamic mechanical analysis (DMA) curve of PE (γ-transition). While it is true that the glass transition displays the most intense peak in the tan δ versus temperature curve, it is also true that in the differential scanning calorimetry (DSC) measurements the ΔCp should be 11.3 J/(K mol) (Wunderlich’s rule) at the glass transition. Gaur and Wunderlich (1980) did not find any heat capacity jump in the 2130
C temperature range, so they insist that the 2128
C DMA peak reflects the freeze of the cistrans equilibrium. But they could not explain why the intensity of the β-transition of PE is very small at 230
C in the DMA recordings. So, it is obvious that there is an onset of some motion at 2128
C. The only question is the type of this motion. Recently, Menczel (2020) applied modulated temperature thermomechanical analysis (MT-TMA) measurements to find the glass transitions in Thermal Analysis of Textiles and Fibers. DOI: https://doi.org/10.1016/B978-0-08-100572-9.00011-2 © 2020 Elsevier Ltd. All rights reserved.

198

Thermal Analysis of Textiles and Fibers

Figure 11.1 Modulated temperature TMA curves of Dyneema ultrahigh molecular mass (UHMM) highly drawn fiber. Solid curve: dimension change; long dashed curve: reversing dimension change; short dashed curve: nonreversing dimension change. The glass transition is clearly seen at 220
C. TA Instruments Q400 TMA. Underlying heating rate 2
C/min; modulation conditions: 6 0.5
C/60 s (Menczel, 2020).

drawn polymeric fibers. In his measurements the reversing dimension change turned out to be much more sensitive in the glass transition measurements than simple traditional DSC curves. He could show the presence of the glass transition in Dyneema fibers at ca. 2 20
C (Fig. 11.1) supporting Gaur and Wunderlich’s conclusions. Thus it is likely that Tg of PE is around 230
C, but the small β-transition in the DMA curves needs some explanation. Polyethylene fibers are very important for practical applications, because polyethylene is an inexpensive polymer. But spinning of PE fibers from the melt is difficult because of the very high melt viscosity of PE. The drawing of melt-spun PE is also difficult and can be done to relatively small draw ratios (up to DR 5 5 3 ) due to the high entanglement density of the PE macromolecules in the melt. Ultraoriented PE fibers were made by Mead et al. (1979) by solid state extrusion. The melting point, crystallinity, density, and thermal shrinkage of these fibers increased with increasing draw ratio. The highest melting point, crystallinity, density, and shrinkage values were obtained at the draw ratio of 2030 3 . The best drawing temperatures for these fibers was 80
C100
C (i.e., somewhat lower than the α-relaxation temperature, which is 110
C for PE). Nevertheless, this technology did not become popular. Ultra-high-molecular-mass polyethylene (UHMMPE, but often having an abbreviation of UHMWPE due to the former expression of “ultrahigh molecular weight”) plays a special role among various PE fibers. It is also called high-modulus polyethylene. Its mass average molecular mass in most cases is between 3.5 and 7.5 million,

Polyethylene fibers

199

but sometimes it is even higher, and the polymer chains are not branched but linear. Thus a very tough material is obtained: the impact strength of this polymer is the highest among all polymers known. The crystallinity of these fibers is very high (usually higher than 85%). Gel-spun fibers from UHMWPE were first commercialized by the Dutch chemical company DSM in 1990. In the gel-spinning manufacturing process of PE fibers, 2% of ultrahigh molecular mass high-density polyethylene (UHMMPE) (with a molecular mass of more than a million) is dissolved in paraffin oil or decaline at around 130
C. A solvent that is used must be a “good solvent” at the dissolution temperature, but PE must crystallize when the solution is cooled. Then spinning is done through relatively large spinnerets (B0.5 mm in diameter). Coming out of the spinneret, the PE solution is cooled, the solvent is removed, and a gel fiber is formed. The solvent removal can be accomplished by extraction of the solvent (wet spinning) or evaporation (dry spinning). After these processes, the gel fiber is drawn at 140
C to very high draw ratios, up to DR . 30 3 , sometimes 130 3 (Russell et al., 2013). The draw ratio depends on both the molecular mass of PE and the concentration of PE in the solution. Higher draw ratios can be achieved with wet drawing than with solvent evaporating dry drawing. These high draw ratios can be achieved, because the macromolecules in the gel are not densely entangled. However, some entanglement is necessary; otherwise the fibers would break during drawing. These fibers are extensively used in areas where very high tensile strength and modulus are needed, like in various defense applications, ballistic protection (bullet proof vests), cut-resistant gloves, various sporting good elements (in helmets for ice hockey, archery, braided rope, protective clothing in the sport of fencing, archery, etc.). They even have space applications. Several Dutch scientists (the research group of Pennings: Smith and Lemstra, 1980; Smook and Pennings, 1984a,b; Hoogsteen et al, 1988) developed the gel-spinning technology of UHMMPE. Gelspun polyethylene fibers became the strongest fibers known to mankind. These fibers have very high orientation and crystallinity. They can be used in the temperature range between 2150
C and 180
C. Their yield strength is around 2.5 GPa, and their Young’s modulus is at or higher than 100 GPa. Due to the small specific density of polyethylene, the strength-to-weight ratio of UHMMPE fibers is about 10 times higher than that of steel. The mechanical properties of UHMMPE fibers exceed those of aramids. Sometimes these fibers are blended with other highperformance fibers, like Vectran, to reduce the creep and stretch. Today, the gel-spun PE fibers are marketed under the trade names of Dyneema (DSM, The Netherlands and the Toyobo/DSM joint venture), Spectra (Honeywell), Tenfor (SNIA Fibre SpA, Italy), and Izanas of Toyobo, Japan. Lately some steps were taken to develop electrospinning of these fibers (Rein et al., 2007). Coming out of the spinneret, the PE solution is cooled, the solvent is removed, and a gel fiber is formed. This can be done by evaporation or by extraction of the solvent. A solvent that is used must be a “good solvent” at the dissolution temperature, but PE must crystallize when the solution is cooled. Smith and Lemstra (1980) and Smook and Pennings (1984b) were the first to apply DSC to study the properties of gel-spun PE fibers. These authors and

200

Thermal Analysis of Textiles and Fibers

Murthy et al. (1990) and Grubb (1983) applied the entanglement concept for development of structural models for high modulus gel-spun PE fibers. The first important paper on thermal properties of ultrahigh-strength gel-spun HDPE fibers was done by Smith and Lemstra in 1980. They prepared the PE gel in paraffin oil solution, and hot drew them at high temperatures (between 90
C and 135
C) up to the draw ratio of 80 3 . The Young’s modulus increased linearly to the draw ratio of 32 3 , but the heat of fusion increased continuously up to the draw ratio of 80 3 , and it was 260 J/g for the highest draw ratio fiber corresponding to 90% crystallinity. Smook and Pennings (1984a,b) performed a more detailed DSC study on morphological and structural changes in hot-drawn UHMMPE fibers. The fibers were gel spun at low-spinning rates (usually 1 m/min) and then were drawn at 148
C in steps. The maximum draw ratio was 80 3 . Due to the slow spinning, the original as-spun fiber consisted mostly of lamellae (not shish-kebabs). At a draw ratio of 6 3 , shish-kebab formation was noticed, while a smooth fibrillary structure could be seen at a draw ratio of 80 3 (Heuvel et al., 1976). 1. Unconstrained melting experiments (Fig. 3.16A): On the DSC curves in unconstrained melting experiments, single (or blurred double), relatively broad (ca.20
C broad) melting peaks were recorded up to the draw ratio of 15 3 , and much narrower (ca.5
C broad) peaks were present for the highest draw ratios (48 3 and 80 3 ). The peak temperature was increasing continuously with the draw ratio from 135
C to 143
C. Similarly, the crystallinity increased from 72% for the as-spun fiber to 93% for the fiber with the draw ratio of 80 3 (the heat of fusion increased from 211 to 274 J/g). It is interesting that the crystallinity of the as-spun fiber was 72%, while it was 52% only for the fiber with DR 5 6 3 (DR indicates the draw ratio). This might have resulted from the fact that the as-spun fiber fully melted when the temperature was raised to the drawing temperature, 145
C. The crystallite size increased strongly from 30 nm for the as-spun fiber to 70 nm for the fiber with DR 5 80 3 (Heuvel et al., 1976). Shrinkage measurements indicated continuously increasing shrinkage from 60
C to the melting point: the value of the shrinkage was 5.5% at 139.5
C. Between 139.5
C and 140
C, a huge, 91% shrinkage took place, and this increased to 97% (145
C). 2. Constrained melting (Fig. 3.16B): The shape of the melting curves was totally different from those described above for the free-to-shrink (unconstrained) melting. In these experiments the recoiling of the chain segments is not as easy as in the free-to-shrink measurements, because the fiber length is kept constant. It is clear from Wunderlich’s experiments (Wunderlich and Czornyj, 1977) that polyethylene under high stresses (e.g., high pressure, or in this case high longitudinal stresses) is often transformed into a pseudohexagonal crystal form (the room temperature, atmospheric pressure crystal form of polyethylene is orthorhombic). This orthorhombic-to-pseudohexagonal phase transformation allowed Wunderlich to prepare the equilibrium PE crystals, because the macromolecular chains could shift more-or-less freely along the z-axis. It is likely that a similar phase transformation took place in Smook and Penning’s experiments.

Smook and Pennings (1984b) compared the heats of fusion values of drawn gelspun polyethylene in the free-to-shrink and constrained type of measurements and observed a significant 7% difference between the two (Figs. 3.163.18). This difference was explained by formation of partially ordered melt in the fixed-length DSC measurements. It is clear that the orientation developed in the melt in the constrained DSC measurements is not stable thermodynamically, and it relaxes slowly,

Polyethylene fibers

201

but the same sample should be melted-crystallized several times before the traces of the original orientation are fully gone. This phenomenon is not unique for polyethylene; it can be observed for other fibers as well. The extent to which this orientation is preserved is different for different fibers. Menczel (2020) concluded that nylons retain this orientation for the longest time in the melt, obviously due to the strong hydrogen bonding. Later, Lemstra et al. (1987) studied the time dependence of drawability of these gel-spun fibers at the drawing temperature of 120
C. They observed the melting peak at 145.5
C at a heating rate of 10
C/min; this clearly indicates superheating of the extended chain crystals. They observed that the Young’s modulus of the fibers increases linearly with increasing draw ratio. Hoogsten et al. (1990) studied by small angle X-ray scattering (SAXS) the influence of the spinning conditions on the drawability, therefore the final mechanical properties of the drawn fibers. Two types of gel-spun fibers were obtained. One of them contained PE crystals of isotropic lamellar structure, while the other one contained shish-kebabs. The shish-kebab containing as-spun fibers could not be drawn to high draw ratios; for that, folded-chain lamellae are needed. Since the excellent mechanical properties of the gel-spun fibers are due to the high draw ratios, spinning must lead to mostly folded-chain lamellae. Wunderlich (2000) showed that multiple melting peaks of highly drawn gel-spun polyethylene fibers originate from sample packing and not the structure of the fiber. Smith and Lemstra (1979) studied how the solvent influences the drawability. The DMA behavior of the α-relaxation was studied in detail by Rong and Williams (1985) on low density polyethylene (LDPE) fibers. This transition occurred at 110
C and as had been shown earlier; it depends most of all on the thickness of the crystallites (Popli et al., 1984). In Rong and Williams’ paper, the β-relaxation took place at 10
C for low-density PE fiber (Fig. 11.2.), but it occurred at a much higher temperature (B55
C) for a commercial high-density PE fiber (Fig. 11.3). Rong and Williams raised the possibility that the β transition could be attributed to the branched structure of low-density polyethylene, but it is not clear why the β-transition occur at such a high temperature in high density polyethylene (HDPE). Rong and Williams (1985) indicated sample problems, and it seems the curves shown in Fig. 11.3 may reflect the DMA behavior of some special PE sample, not the usual HDPE fiber. However, the DSC measurements made by Gaur and Wunderlich (1980) and the MT-TMA measurements made by Menczel (2020) indicate that the β-transition is likely to be the glass transition of PE. Rong and Williams (1985) obtained the γ-transition at 290
C for both LDPE fiber and HDPE commercial fiber. Despite of Beatty and Karasz’ (1979) opinion, this relaxation may correspond to the concerted motion of several methylene units in the chain [crankshaft motion of Shatzki (1962)]. Since the crankshaft mechanism of Shatzki indicates the onset of segmental motion of several beads (independent mobile units), it is possible that the glass transition of polyethylene takes place in two steps (at 290
C and 220
C): at 290
C a “microglass transition” takes place (extending to 45 CH2-units), while the full unfreezing occurs at 220
C.

202

Thermal Analysis of Textiles and Fibers

Figure 11.2 The tan δ versus temperature DMA curves for a low-density PE monofilament fiber drawn to various ratios. The draw ratios were the following: curve A, DR 5 2.0; curve B, DR 5 2.7; curve C, DR 5 3.5. Source: From Rong, S.-D., Williams, L.H., 1985 J. Appl. Polym. Sci. 30, 2575. Rheovibron viscoelastometer, frequency 110 Hz. Reprinted with permission of Wiley.

Figure 11.3 The tan δ versus temperature DMA curves for a commercial HDPE monofilament fiber at various frequencies. The frequencies were the following: curve A, 11 Hz; curve B, 35 Hz; curve C 110 Hz. The α relaxation at ca.110
C could not be found in these experiments (not shown in the figure). Intensive β relaxation implies that the fiber may contain a large number of chain branches. As expected, the γ relaxation can be seen at 2100
C. Source: From Rong, S.-D., Williams, L.H., 1985. J. Appl. Polym. Sci. 30, 2575, Figs. 4 and 8. Rheovibron viscoelastometer, frequency 110 Hz. Reprinted with permission of Wiley.

Polyethylene fibers

203

Murthy et al. (1990) applied simultaneous DSCX-ray diffraction to study the melting behavior of highly oriented PE fibers. Smook and Pennings (1984b) ascribed the middle endothermic peak in Fig. 3.16B to the orthorombic!pseudohexagonal phase transition, and this conclusion was supported by the experiments of Murthy et al. (1990). They recorded the highest temperature melting peak in the constrained state experiments at 160
C162
C (20
C higher than the equilibrium melting point of PE!, an excellent proof for superheating).

References Beatty, C.L., Karasz, F.E., 1979. J. Macromol. Sci. Rev. Macromol. Chem. C17, 37. Gaur, U., Wunderlich, B., 1980. Macromolecules 13, 445. Grubb, D.T., 1983. J. Polym. Sci. Polym. Phys. Ed. 21, 165. Heuvel, H.M., Huisman, R., Lind, K.C.J.B., 1976. J. Polym. Sci. Polym. Phys. Ed. 14, 921. Hoogsteen, W., Kormelink, H., Eshuis, G., Ten Brinke, G., Pennings, A.J., 1988. J. Mater. Sci. 23, 3467. Hoogsten, W., Pennings, A.J., ten Brinke, G., 1990. Colloid. Polym. Sci. 268, 245. Lemstra, P.J., van Aerle, N.A.J.M., Bastiansen, C.W.M., 1987. Polym. J. 19, 85. Mead, W.T., Desper, C.R., Porter, R.S., 1979. J. Polym. Sci. Polym. Phys. Ed. 17, 859. Menczel, J.D., 2020. To be published. Murthy, N.S., Correale, S.T., Kavesh, S., 1990. Polym. Commun. 31, 50. Popli, R., Glotin, M., Mandelkern, L., 1984. J. Polym. Sci. Polym. Phys. Ed. 22, 407. Prime, R.B., Wunderlich, B., 1969. J. Polym. Sci. Polym. Phys. Ed. 7, 2061. Prime, R.B., Wunderlich, B., Melillo, L., 1969. J. Polym. Sci. Polym. Phys. Ed. 7, 2091. Rein, D.M., Shavit-Hadar, L., Khalfin, R.L., Cohen, Y., Shuster, K., Zussman, E., 2007. J. Polym. Sci. Part B: Polym. Phys. 45, 766. Rong, S.-D., Williams, L.H., 1985. J. Appl. Polym. Sci. 30, 2575. Russell, B.P., Kandan, K., Deshpande, K.V.S., Fleck, N.A., 2013. Int. J. Impact Eng. 60, 1. Shatzki, T.F., 1962. J. Polym. Sci. 57, 496. Smith, P., Lemstra, P.J., 1979. Makromol. Chem. 180 (12), 2983. Smith, P., Lemstra, P.J., 1980. J. Mater. Sci. 15 (2), 505. Smook, J., Pennings, J., 1984a. J. Mater. Sci. 19, 31. Smook, J., Pennings, J., 1984b. Colloid. Polym. Sci. 262, 712. Wunderlich, B., 2000. Private communication to Menczel, J. Wunderlich, B., Czornyj, G., 1977. Macromolecules 10, 905.

Further reading Boller, A., Wunderlich, B., 1996. 11th Congress of Thermal Analysis and Calorimetry, PA, August 1216. Lemstra, P.J., Smith, P., 1979. US Patent US4430383 A, Priority Date Jun 27, 1979, Filing Date Sep. 30, 1982. Ziabicki, A., 1976. Fundamentals of Fiber Formation. John Wiley and Sons, New York.

Polypropylene fibers

12

Alfre´d Menyha´rd1, Joseph D. Menczel2 and Tonson Abraham3 1 Laboratory of Plastics and Rubber Technology, Department of Physical Chemistry and Materials Science, Budapest University of Technology and Economics, Budapest, Hungary, 2 Thermal Measurements LLC, Fort Worth, TX, United States, 3Retired from ExxonMobil Chemical Company, Avon, OH, United States

Abstract Polypropylene (PP) is an olefin polymer (so it is thermoplastic) with a relatively low melting point (B170 C). It is important industrially, because it is used in a number of applications. PP fibers started became popular in the 1970s. The 2014 worldwide production value was US$56.73 million. This means that PP fiber is the fourth largest volume artificial fiber after the polyesters, nylons, and acrylics. The difference from these three fibers is that PP fiber is hardly ever used as apparel, and its major use is for industrial applications: carpets, geotextiles, ropes and reinforcing fiber in concrete and soil, sanitary products, and surgical sutures. A lot of information on PP resin is described in the Polypropylene Handbook.

12.1

Introduction

Polypropylene (PP) is an olefin polymer (so it is thermoplastic) with a relatively low melting point (B170 C). It is important industrially, because it is used in a number of applications. PP fibers became popular in the 1970s. The 2014 worldwide production value was US$56.73 million (Crystal Market Research, 2017). This means that PP fiber is the fourth largest volume artificial fiber after polyesters, nylons, and acrylics. The difference from these three fibers is that PP fiber is hardly ever used as apparel, and its major use is for industrial applications: carpets, geotextiles, ropes and reinforcing fiber in concrete and soil, sanitary products, and surgical sutures. A lot of information on PP resin is described in the Polypropylene Handbook (Pasquini, 2005). PP is a stereoregular polymer, which can be isotactic, syndiotactic, or atactic. When PP is synthesized it always has some atactic content. At the beginning of the PP era, this atactic PP was removed from just synthetized PP by solvents. Now, with much better catalysts (metallocene) the stereoregularity is controlled much better. Isotacticity of at least 95% is needed for fiber production. PP chain conformation is helical (31-helix), it is a semicrystalline polymer, with several crystal forms: α-form (monoclinic), β-form (pseudohexagonal), and the γ-form (orthorhombic). Another crystal form, called the ε-form, has been recently found, but it is very rare and can be formed in the stereodefective parts of the Thermal Analysis of Textiles and Fibers. DOI: https://doi.org/10.1016/B978-0-08-100572-9.00012-4 © 2020 Elsevier Ltd. All rights reserved.

206

Thermal Analysis of Textiles and Fibers

sample (Lotz, 2014). The most common crystal form is the α-form, this is what develops in melt crystallization of PP. The β- and the γ-crystal forms are the result of stress crystallization or crystallization induced by nucleating agents. The α-form is the most important commercially. The β-form sometimes is used due to its somewhat higher impact resistance. However, the impact resistance can be increased by copolymerization with a small amount of ethylene comonomer (7 series ProFax, the former Hercules, presently LyondellBasell). In addition, there is a somewhat ordered state of PP called the “mesophase.” It is unknown whether it is some imperfect crystal form, but it is definitely not a liquid crystalline phase despite of earlier suggestions (see below). Application of PP fibers is important in areas, such as the manufacturing of ropes, concrete reinforcement additive (geotextiles, reinforcing fiber) to reduce cracking or spalling, increase soil strength; as sutures and reinforcing meshes for hernia and female incontinence repair (see Chapter 15: Surgical sutures), baby diapers, sanitary products (nonstabilized PP); and filters for filtering liquids and various gases. PP fiber is not really used as apparel, but as a nonwoven, it has major applications in home furnishings (carpet backing and furniture fabrics) and automotive markets. It is extensively used for wall covering, in warm-weather clothing, as surgical disposable fabric. PP fibers are distinguished as having the smallest thermal conductivity among fibers: K 5 0.11 W/(m C) (K is the thermal conductivity), much lower than for Nylon 6 (0.25) or PMMA (0.21), therefore PP fibers have great insulation properties. Another major advantage of PP fiber is its price. But its melting point is relatively low (170 C), and it is not easy to texturize or dye it: there are no polar groups of the fiber surfaces, thus the fiber does not really hold any other material on its surface. In the future, plasma treatment of the fiber surface may be the solution to this problem: with plasma treatment polar groups can be attached to the surface that would be able to hold the dye molecules. PP has excellent resistance to chemicals (except oxidizing agents). It has low moisture absorption (less than 0.1%), therefore its dimensional stability in humid environment is very good. It is not soluble in most solvents, that is why gel spinning would be very expensive for this fiber (it can be dissolved in some solvents at temperatures close to its melting point). PP fiber has extremely low density having less than 1. PP is the lightest of all synthetic fibers, its heat conductivity is the smallest among all fibers, and it is an excellent insulator in both heat and electricity. It is combustible but not easily ignitable and also cannot be considered highly flammable. With incorporating certain additives, it can be made flame-retardant. A serious issue of PP fibers is its poor thermal, thermo-oxidative, and light stability. In order to improve these characteristics (coming from the presence of the tertiary carbon atom in the chain), significant amounts of stabilizers need to be added. This limits its application as a food container or medical applications. The only way to bypass these issues is not to add stabilizers to the material, but then it can be used for a short time only (medical grade PPs). This certainly does not limit its application as a surgical suture. When PP used at room temperature or higher temperature for a long time, significant creep can be observed. Nonwovens are one of the major applications of PP fibers. Nonwovens are web structures: in them the filaments are bonded thermally (partial melting), chemically,

Polypropylene fibers

207

or mechanically (e.g., baby diapers and automobile industry applications). PP nonwovens are particularly well suited for manufacturing hygienic products, for example, baby diapers, but for these applications special PP is used without stabilizing agents (the so-called medical grade PP). The mechanical properties of PP fibers are diverse depending on the specific applications. In most cases these mechanical properties are not really good, but in the fields of their use is no need for very good mechanical properties. For example, for use as sutures, no superhigh tenacities are required. The tenacity of PP fibers is usually 4.0 5.5 cN/ dTex. For use in ropes or nets, yarns can be used with higher tenacities of up to 8.0 cN/ dTex. The mechanical properties of fibers are mostly determined by chain orientation that is controlled by the spinning and drawing conditions. For manufacturing high tenacity PP fibers, special techniques are used, such as ultradrawing (Weeks and Porter, 1975), solid-state extrusion (Lim et al., 1989), and crystal surface growth (Baumann, 1963). Tenacities of up to 11.5 cN/dTex can be achieved, but these are rare products. Another application area is geotextiles (reinforcing fibers in concretes, road materials, or soil), especially in areas where photodegradation is not a serious issue. The disadvantage of PP fibers in these applications is temperature: even at room temperature PP fibers tend to creep (Tg of PP is relatively low). Hollow fibers are also made of PP. As their name says, hollow fibers are small microporous tube-like fibers having diameter around 200 μm. These fibers are extensively used for ultra- and micro-filtration purposes (also in waste water treatments), in the chemical industry, and medical applications. Hollow fiber membranes are extensively used for separating dissolved gases from liquids, for example, water (Heo and Park, 2014). Of course, there are some negatives for PP fibers. As mentioned earlier, due to the tertiary carbon atom in the main chain, PP is easily susceptible to thermal, thermo-oxidative, and photodegradation. Its oxidative degradation increases under stresses (problem in the geotextile area).

12.2

Manufacturing of polypropylene fibers

The commercial PP fibers are made by traditional melt spinning. The schematic of melt spinning is shown in Fig. 12.1 and the drawing is in Fig. 12.2.

12.3

Processing of polypropylene

The most important thermal analysis parameter of PP during processing is the melting point. For most PPs, this is around 170 C, but it is lower for the popular PP copolymers with low ethylene comonomer content (e.g., the ProFax 7-series polymers) because the PP crystallites in these polymers are considerably smaller than in the homopolymer PP: the short ethylene chain sequences in the chain will limit the crystal size. Melt temperatures in injection molding are between 220 C and 250 C, and the mold temperature is between room temperature and 65 C. The mold temperature

208

Thermal Analysis of Textiles and Fibers

Melt spinning Feed

Polymer chips

Melting Metered extrusion Filter and spinnaret Cooling and solidifying

Lubrication

Feed rolls

Yarn driving Packaging

Bobbin

Packaging

Bobbin drive roll

Figure 12.1 Schematic of melt spinning (Billmeyer, 1984). Source: From Billmeyer, F.W., 1984. Textbook of Polymer Science. Wiley (Interscience), New York and London, p. 493, Fig. 18.4/a. Reprinted with permission of Wiley.

will determine the cooling rate of the polymer, so the final crystallinity of the product. In extrusion the temperature between the throat and the front is between 200 C and 230 C, but for cast films and fibers this temperature is higher. It is a very tricky issue to select the melt temperature in PP processing. High melt temperatures will cause considerable degradation (so most PP products contain stabilizers) but will lower the melt viscosity. At low melt temperatures the opposite is observed. In fiber manufacturing the melt temperatures are somewhat higher than in injection molding, so the product is even more susceptible to degradation. Roberts (1982) developed a PP filament extrusion with melt temperatures as low as 177 C, but such a processing did not become popular (probably due to high melt viscosity).

12.4

Properties of polypropylene fibers

PP fibers, like all semicrystalline fibers, are composed of a crystalline fraction and one or two or three amorphous regions (mobile amorphous fraction, rigid amorphous fraction, and oriented amorphous fraction).

Polypropylene fibers

209

Figure 12.2 Schematic of fiber drawing: V1 , V2, drawing; V1 . V2, relaxation (Billmeyer, 1984). Source: From Billmeyer, F.W., 1984. Textbook of Polymer Science. Wiley (Interscience), New York and London, p. 494, Fig. 18.5 Reprinted with permission of Wiley.

The diameter of the spherulites in the as-spun fibers is between a micrometer and several millimeters. Each crystal is surrounded by noncrystalline, amorphous material. Fiber spinning and drawing will orient both the crystalline and the amorphous regions. When the as-spun fiber is stretched, the spherulite deformation at small extension levels is elastic, and the supermolecular structure will not change. However, the spherulites become highly oriented at large deformations, and at very large deformation levels the spherulites are changed into microfibrils imparting highly anisotropic properties to the fibers. During the industrial processing, definite crystallization takes place. PP in this aspect is not similar to polyesters. In addition to the presence of the crystals and the amorphous phases (mobile, rigid, and oriented), a “mesophase” exists in quenched PP samples (the optimum rate of cooling for obtaining the mesophase is between 100 C/s and 1000 C/s). The formation of the mesophase takes place between 0 C and 80 C (Natta and Corradini, 1960; Natta, 1960). It is almost certain that this mesophase is not a liquid crystal (Androsch et al., 2010). The mesophase preferably forms in quenched PPs, and in temperature-resolved wide angle X-ray diffraction (WAXD) experiments it was shown to have a broad double diffractogram (Nishida et al., 2012). Silvestre et al. (2007) proved the existence of bimodal crystallization for PP by fast-scan differential scanning calorimetry (FS-DSC): in the 60 C 120 C temperature range, the development of the monoclinic crystal form (α-) takes place, while the mesophase formation was observed between 10 C and 60 C (that is

210

Thermal Analysis of Textiles and Fibers

between Tg and 60 C). According to Schick (Schick et al., 2017; Schick, 2018), the bimodal crystallization is seen for many polymers, also for polymers without the mesophase formation. It is believed this bimodal crystallization is the consequence of homogeneous nucleation at low temperatures and heterogeneous nucleation at higher temperatures. The nucleation density in homogeneous nucleation is often extremely high, particularly for isotactic polypropylene (iPP). It has several orders of magnitude higher than the nucleation density in heterogeneous nucleation. The homogeneous nucleation density can be so high that the space for crystal growth is extremely limited, so the crystallization growth is terminated at the start of the crystallization process, and with a large amount of the rigid amorphous fraction. This may favor the formation of the mesophase instead of the α-crystal form. According to these arguments, the mesophase must be something very close to the rigid amorphous phase. On heating the mesophase changes to the α-form (nuclear magnetic resonance (NMR)) measurements, see Schaefer et al., 1990). It is questionable whether the helical structure of PP plays some role in the mesophase formation. It is not clear whether the easy raw nucleation of PP can be connected somehow to the mesophase, since both these may have common roots with the PP helix. It needs to be mentioned that according to X-ray measurements, the lateral correlation of the chain segments in PP is closer to the monoclinic form than to the pseudohexagonal (which is the β-form) structure. Therefore the mesophase cannot be explained as a “disrupted crystal phase” similar to the pseudohexagonal phase of polyethylene. The mesophase has been observed in fibers taken up at low speed, rapidly cooled in water with the addition of ice, extruded at high temperatures, and extruded from PP with low molecular mass. Thus PP probably has three crystal forms: the α-form is monoclinic, and the β-form has pseudohexagonal structure, and the γ-form is orthorombic (Padden and Keith, 1959; Keith et al., 1959). Use of certain nucleating agents promotes the formation of the β-form; the presence of chain defects (disruption of isotacticity) and low molecular mass fractions or crystallization at elevated pressures can induce development of the γ-form (orthorhombic unit cell) (Sauer and Pae, 1968; TurnerJones, 1971). In the fibers mostly α-crystals are present. It is the most stable crystal form of PP. However, fibers containing the β-form can also be manufactured. The structure of the mesophase is still debated in the literature (Nishida et al., 2012). Commercial PP fibers are made by melt spinning with subsequent hot (or sometimes cold) drawing. High-modulus, high-strength PP fibers can also be produced. There are several techniques of doing that, which are as follows: (1) high mesophase content PP fiber can be made by quenching, and then this fiber can be drawn at low rates to very high draw ratios at temperatures close to the melting point (130 C 135 C) (Sheehan and Cole, 1964; Nadella et al., 1977). (2) Clark and Scott (1974) developed a two-stage drawing process at specific controlled temperatures. (3) Kunugi applied his zone-drawing zone-annealing technique to PP (Kunugi et al., 1983). However, these fibers are not commercial. They can be adequate for specific applications but are not suitable for mass production. As explained earlier, the cooling rate plays an important role in PP fiber production. Take-up velocity is an important parameter in the fiber processing, changing of which can drastically change the cooling rate, and this will change the

Polypropylene fibers

211

crystallization temperature, thus the mesophase content. At very high take-up velocities the stress may increase considerably due to the drastic increase of the chain orientation. At high molecular orientation the crystallization temperature increases because the oriented regions have nucleating effect and at the same time the mechanism of the crystallization process changes: bulk crystallization slowly changes to oriented crystallization.

12.5

Thermal analysis of polypropylene fibers

12.5.1 Melting PP is an olefin polymer. Like all other crystalline polymers, it is semicrystalline. Minimum 95% isotacticity is required to spin PP fibers. The equilibrium melting point of the monoclinic (α-crystal form) of isotactic PP (Tmo ) is still debated in the literature. Unlike polyethylene, extended-chain macroscopic PP crystals could never be prepared, therefore the equilibrium melting point is unknown, it was estimated from the measurements on metastable crystals. From these measurements, Tmo of PP is between 187.5 C and 220 C (Cox and Duswalt, 1967; Fatou, 1971; Samuels, 1975a,b; Gaur and Wunderlich, 1981; Wunderlich, 1980). The most likely Tmo value is 220 C. Samuels estimated Tmo of the β-crystal form (pseudohexagonal crystal form) to be 170 C (Samuels, 1975a). The experimental nonequilibrium melting point of PP (mostly α-crystal form) is around 170 C. The glass transition temperature of isotactic PP originally was reported to be 214 C (Passaglia and Kevorkian, 1963; Gaur and Wunderlich, 1981). Since the use of metallocene catalysts, the isotacticity of the manufactured PP is higher than before. The crystallinity is also increased. This may be the reason why the glass transition temperature of new PP variants is around 25 C to 12 C (Menczel, 2020). The peak temperature of melting (Tmp) and the melting point (Tm, the highest temperature (last point) of the melting endotherm) of PP are usually around of 165 C and 170 C. The heat of fusion of 100% α-crystalline PP was estimated to be 165 J/g. During cooling the crystallization of iPP starts at 125 C 110 C, but the crystallization temperature can be manipulated considerably using efficient nucleating agents (Menczel and Varga, 1983; Fairgrieve, 2005; Mitsuishi, 1996; Thierry et al., 1992). The most important technique for improving the impact resistance of iPP is copolymerization of propylene with a small amount of ethylene comonomers (Galli et al., 1995; Albizzati et al., 1996), but the fibers containing the β-form have also higher impact strength. The presence of ethylene comonomers may interfere with the regularity of the polymer chains, and the characteristic values of melting and crystallization change considerably. The crystallinity as well as heat of fusion decreases in random copolymers. The melting and crystallization temperatures are also lower. As previously mentioned, iPP is a polymorphic material with three crystal forms. However, two crystal forms only (α- and β-) have commercial importance. The β-form melts around 150 C, and its heat of fusion is smaller than for the α-form (naturally, it is pseudohexagonal). The advantage of the β-form is its large impact resistance (twice as much as the impact resistance of the α-form). It can be

212

Thermal Analysis of Textiles and Fibers

Figure 12.3 Melting peaks of drawn PP fibers. The draw ratio for the samples as-spun, 1, 2, 3, 4, 5, and 6 was 1 (undrawn), 1.39, 1.84, 2.29, 2.70, 3.21, and 3.58, respectively. The peak temperatures are the following: as-spun fiber 163.7 C; (1) 166.3 C; (2) 164.8 C and 165.5 C; (3) 168.0 C and 181.0 C; (4) 170.0 C, 172.9 C, and 187.3 C; (5) 173.9 C and 192.0 C; (6) 172.4 C and 193.8 C. Source: From Samuels, R.J., 1975a. J. Polym. Sci., Polym. Phys. Ed. 13, 1417, Fig. 1. Reprinted with permission from Wiley.

manufactured by adding efficient nucleating agents (Varga, 1995; Thierry et al., 1992; Libster et al., 2007; Menyha´rd et al., 2006). The α-form is enantiotropic, while the β-form seems to be monotropic. The first detailed thermal analysis measurements on PP fibers were done by Samuels (1970; 1975a,b; 1979). He used as-spun fibers, hot drawn fibers, and heatset fibers to study the melting behavior of PP fibers in free-to-shrink and constrained modes. The draw ratio of these fibers was between 1.4 and 4.5. The melting curves of drawn PP fibers in constrained state are shown in Fig. 12.3. Samuels used three parameters for structural characterization of the fibers: (1) crystallinity (X); (2) orientation function of the chains in the crystalline phase (fX), and (3) orientation function of the chains in the amorphous phase ( fAV). As usual for semicrystalline fibers, the thermal and mechanical history of the fiber determines the melting properties. The melting properties of fibers (and polymers in general) are determined by the thermal and mechanical history of the sample, but the state of the sample will also play some role. Several scientists noticed that different melting curves are recorded for fibers when they are kept at a constant length or fibers that can shrink during the melting (Miller, 1971; Mead and Porter, 1976; Miyagi and Wunderlich, 1972; Samuels, 1975a,b; Tashiro et al., 1980; Todoki and Kawaguchi, 1977a,b; see Chapter 3: Differential scanning calorimetry in fiber

Polypropylene fibers

213

Figure 12.4 The peak temperature of melting of the higher temperature melting peak (T2M) as a function of amorphous orientation function (fAV) of the drawn PP fibers (Samuels 1975a). Source: From Samuels, R.J., 1975a. J. Polym. Sci., Polym. Phys. Ed. 13, 1417, Fig. 3. Reprinted with permission from Wiley.

research). In Samuels’ experiments the melting point of PP was significantly higher in the constrained measurements than in the free-to-shrink measurements, as is the case for almost all drawn fibers. Samuels found two melting peaks for the fibers when recorded at fixed length. He noticed that the peak temperature of melting for these two peaks did not depend on the crystal orientation but showed considerable dependence on the amorphous orientation function (Fig. 12.3). The melting point values were extrapolated to 220 C for full amorphous orientation ( fAV 5 1.0, see Fig. 12.4). Thus Samuels made a conclusion that Tmo of the monoclinic crystals is 220 C. There was another, lower temperature melting peak on the DSC traces, but the author could not explain the origin of this melting peak. It is possible that the constraining procedure during the sample preparation was not sufficient, and the lower temperature melting peak may reflect the melting of shrinkable parts of the fiber specimens: the lower Tmp of the drawn fibers have similar value as Tmp of the as-spun fiber. Jaffe (1978) made a series of measurements on melting and crystallization of asspun PP fibers. He received increasing melting point with decreasing heating rate when the spinning stress was increasing. This could mean reorganization during melting. The time dependence of the melting indicated that two crystalline morphologies, fibrillar and lamellar, develop when the stress level during spinning is higher than a certain value. His results indicated that the morphology shifted toward one containing less fibrillar and more spherulitic structure as the melt temperature during spinning was increasing.

Figure 12.5 Comparison of melting and crystallization curves of drawn PP fibers run in freeto-shrink and constrained modes. (A) Melting curves recorded in the constrained and free-toshrink experiment. (B) The melting curves in constrained state: first heating: as-received fiber (solid curve), second heating: melting after crystallizing at a rate of 10 C/min (dashed curve); the final temperature of heating was 220 C. TA Instruments Q2000 DSC, heating rate 5 10 C/min. (C) The crystallization in constrained and free-to-shrink measurements, cooling rate 5 10 C/min (solid line: constrained; dashed line: free-to-shrink). (Menczel, 2020).

Polypropylene fibers

215

Menczel (2020) carried out a series of experiments with traditional and modulated temperature DSC studying the reorganization process during melting of samples in free-to-shrink and constrained modes. The melting point was considerably higher in the constrained mode measurements than in the free-to-shrink experiments (Fig. 12.5A). Similar crystallization curves were obtained when the fibers were melted in constrained and free-to-shrink modes (Fig. 12.5B). Finally, when these samples were reheated, the melting was almost identical for the constrained and free-to-shrink samples. Since the samples in the first heating were heated to 220 C, this indicates that the raw nuclei were completely destroyed at this melt temperature. First time, MT-DSC measurements were compared for fiber meltings carried out in constrained and free-to-shrink modes (Fig. 12.6). When the experiments are done in constrained (fixed length) mode, essentially no reorganization takes place at the start of the melting process as indicated by the non-reversing heat flow curve (Fig. 12.6A), but when the measurement is done in free-to-shrink mode (fiber pieces chopped up to small pieces to allow shrinkage), the non-reversing heat flow curve shows significant reorganization during melting (Fig. 12.6B).

12.5.2 The glass transition It is not easy to record the glass transition of PP fibers by conventional DSC. As customary for semicrystalline polymers, especially for drawn fibers, the glass transition is shallow and extremely broad due to the presence of the rigid (Menczel and Wunderlich, 1981) and the oriented amorphous fractions. The glass transition of the as-spun PP fiber is shown in Fig. 12.7. The enthalpy relaxation (hysteresis peak) in the DSC curves at the glass transition of amorphous polymers is a usual phenomenon: when the sample is reheated at a faster rate than it had been previously cooled, an endothermic hysteresis peak is overlaid on the heat capacity jump. However, Menczel and Wunderlich (1981) noticed that the hysteresis peak is missing for semicrystalline polymers (conventional DSC!), and they tried to connect this phenomenon to the formation of the rigid amorphous phase. Later Menczel (2020) noticed that the enthalpy relaxation can be seen in the nonreversing heat flow when these experiments are recorded in modulated mode (Fig. 12.8). Since it is very difficult to see the glass transition of drawn fibers with conventional DSC, this phenomenon can be used to determine the glass transition. Menczel (2020) also obtained indication that the glass transition temperature is higher when the orientation of the sample is higher (Fig. 12.9). When the fiber is run in the constrained mode, melted, and rerun after crystallization, Tg decreases somewhat. Menczel (2020) could record the glass transition with modulated temperature thermomechanical analysis (MT-TMA) (Fig. 12.10). In this case, he could even separate the glass transition of the mobile amorphous fraction and the glass transition of the rigid amorphous fraction that was not possible before.

216

Thermal Analysis of Textiles and Fibers

Figure 12.6 (A) Melting curves of drawn PP fiber in the constrained (fixed-length) mode recorded by MT-DSC. Some weak reorganization can be seen in the nonreversing heat flow signal at around 155 C. The curves from top to bottom: reversing heat flow, total heat flow, non-reversing heat flow (B) In the free-to-shrink mode the nonreversing heat flow indicates significant reorganization during melting. The curves from top to bottom: reversing heat flow, total heat flow, non-reversing heat flow. TA Instruments Q2000 DSC, underlying heating rate 5 C/min. Modulation conditions: 6 0.5 C/min/40 s (Menczel, 2020).

12.5.3 Comparison of results/experimental conditions When running any fiber by DSC, including PP, it is important to pick up similar samples and to prepare the samples in the same way. Otherwise the runs cannot be compared. First, it needs to be decided whether free-to-shrink or constrained

Polypropylene fibers

217

Figure 12.7 The glass transition of as-spun iPP fiber as recorded by conventional DSC. Menczel, unpublished results. TA Instruments Q2000 DSC, heating rate 10 C/min.

Figure 12.8 The hysteresis peak of drawn PP fiber in the nonreversing heat flow signal recorded in the free-to-shrink mode at a rate of heating of 10 C/min after the samples were cooled at the rates of cooling of 1 C/min, 2 C/min, 5 C/min, and 10 C/min from 50 C. TA Instruments Q2000 DSC. Underlying heating rate is 10 C/min. Modulation conditions: 6 0.5 C/40 s (Menczel, 2020).

measurements should be done. Free-to-shrink measurements are usually simpler. The fiber must be chopped to short pieces and run in standard DSC pans. Care must be taken to avoid knots and bundles. In such sample parts the shrinkage may be vastly different, and the shape of the DSC curve will be altered (see Fig. 12.11).

218

Thermal Analysis of Textiles and Fibers

Figure 12.9 The glass transition of polypropylene fiber when the measurement is carried out in the constrained (fixed-length) mode. The sample was melted, then crystallized at a cooling rate of 10 C/min, and reheated at a rate of 10 C/min. The glass transition during the first heating (sample with higher orientation) is higher than in the second heating (orientation effect). TA Instruments Q2000 DSC, heating rate 10 C/min (Menczel, 2020).

Figure 12.10 The glass transition of drawn PP fiber recorded with MT-TMA. The glass transition of the mobile amorphous phase [Tg(mob)] and the glass transition of the rigid amorphous phase [Tg(rig)] are clearly visible on the reversing dimension change signal (Menczel, 2020). Tg(mob) is from 220 C to 28 C (midpoint 5 214 C), Tg(rig) is from 28 C to 125 C (midpoint 5 9.0 C). Underlying heating rate 2 C/min. Modulation conditions: 6 0.5 C/60 s (Menczel, 2020).

Polypropylene fibers

219

The mass of the samples must be similar. In most cases 1 mg would be the optimum sample mass because at this mass the resolution of the instrument is the best and the thermal lag is not significant (see Fig. 12.12). The heating rate must also be standardized. Although very slow heating would be preferred to increase the temperature resolution, there is no time to make extremely slow runs in the industry. Therefore 5 C/min or 10 C/min is recommended.

Figure 12.11 Small fiber pieces of polypropylene. Effect of the various sample pieces on the shape of the DSC curve. Endo down. Source: From Steinmann, W., Walter, S., Beckers, M., Seide, G., Gries, T., 2013. Chapter 12: Thermal analysis of phase transition in polymeric fibers. In: Elkordy, A.A. (Ed.), Applications of Calorimetry in a Wide Context, Differential Scanning Claorimetry, Isothermal Titration Calorimetry and Microcalorimetry. IntechOpen Fig. 5a. Courtesy of IntechOpen.

Figure 12.12 Polypropylene small fiber pieces. Effect of the sample mass on the resolution, heating rate 5 10 C/min. Endo down. Source: From Steinmann, W., Walter, S., Beckers, M., Seide, G., Gries, T., 2013. Chapter 12: Thermal analysis of phase transition in polymeric fibers. In: Elkordy, A.A. (Ed.), Applications of Calorimetry in a Wide Context, Differential Scanning Claorimetry, Isothermal Titration Calorimetry and Microcalorimetry. IntechOpen Fig. 6, Courtesy of IntechOpen.

220

12.6

Thermal Analysis of Textiles and Fibers

Conclusion

PP fiber is an important commercial fiber. This fiber is the fourth largest volume commercial fiber on the world market, and its major advantage is that it is cheap and can be used in many different areas. It is not used as apparel but very popular in nonwoven applications. It can be expected that the use of this fiber will be increasing in certain areas since the crystallinity and the isotacticity of PP is higher than used to be due to new, more efficient metallocene catalysts. Obviously, the development in the polymerization process of propylene is not over, and the goal is to achieve 100% isotacticity. When this is done, the mechanical properties of PP fiber will further improve.

References Albizzati, E., Giannini, U., Collina, G., Noristi, L., Resconi, L., 1996. Catalysts and Polymerizations. In: Moore, E.P. (Ed.), Polypropylene Handbook. Hanser/Gardner Publications, Cincinnati, OH, p. 11. Androsch, R., Di Lorenzo, M.L., Schick, C., Wunderlich, B., 2010. Polymer. (Guildf) 51 (21), 4639. Baumann, H.P., 1963. Am. Dyestuff Rep. 52, 527. Billmeyer, F.W., 1984. Textbook of Polymer Science. Wiley (Interscience), New York and London. Clark, E.S., Scott, L.S., 1974. Polym. Eng. Sci. 14 (10), 682. Cox, W.W., Duswalt, A.A., 1967. Polym. Eng. Sci. 7, 1. Crystal Market Research, Research Study: Polypropylene Fiber Market By Process, Form and End User-Global Industry Analysis and Forecast to 2023. ,globenewswire.com/ news-release/2017/12/06/1234107/0/en/Polypropylene-Fiber-Market-By-Global-IndustryAnalysis.. Fairgrieve, S., 2005. Rapra Rev. Rep. 16, 1. Fatou, J.G., Eur. Polym. J. 7, 1971, 1057. Galli, P., Haylock, J.C., Simonazzi, T., 1995. Manufacturing and Properties of Polypropylene Copolymers. In: Karger-Kocsis, J. (Ed.), Polypropylene: Structure, Blends and Composites, Vol. 2. Chapman & Hall, London, p. 1. Gaur, U., Wunderlich, B., 1981. J. Phys. Chem. Ref. Data, 10, 1051. Heo, P.W., Park, I.S., 2014. Int. J. Agric. Biosyst. Eng. 8 (7), 1295. Jaffe, M., 1978. Thermal Methods in Polymer Analysis, (ed by Shalaby, S.W.), p. 93, Franklin Inst. Press, Philadelphia, PA. Keith, H.D., Padden Jr., F.J., Walter, N.M., Wyckoff, H.W., 1959. J. Appl. Phys. 30, 1485 1488. Kunugi, T., Ito, T., Hashimoto, M., Ooishi, M., 1983. J. Appl. Polym. Sci. 28, 179. Libster, D., Aserin, A., Garti, N., 2007. Polym. Adv. Technol. 18 (9), 1. Lim, J.G., Gupta, B.S., Waller, G., 1989. Progr. Polym. Sci. 14, 763. Lotz, B., 2014. Macromolecules 47 (21), 7612. Mead, W.T., Porter, R.S., 1976. J. Appl. Phys. 47 (C10), 4278. Menczel, J., Varga, J., 1983. J. Therm. Anal. 28, 161.

Polypropylene fibers

221

Menczel, J., Wunderlich, B., 1981. J. Polym. Sci., Polym. Lett. Ed. 19 (5), 261. Menczel, J.D., 2020 to be published. Menyha´rd, A., Varga, J., Molna´r, G., 2006. J. Therm. Anal. Cal. 83 (3), 625 630 (2006). Miller, B., 1971. Themochim. Acta 2 (3), 225. Mitsuishi, K., 1996. Polypropylene, nucleating agents. In: Salamone, J.C. (Ed.), The Polymeric Materials Encyclopedy, vol. 9. CRC Press, Boca Raton, FL, p. 6602. Miyagi, A., Wunderlich, B., 1972. J. Polym. Sci., Polym. Phys. Ed. 10, 1401. Nadella, H.P., Heoson, H.M., Spruiell, J.E., White, J.L., 1977. J. Appl. Polym. Sci. 21, 3003. Natta, G., 1960. Makromol. Chem. 35, 94. Natta, G., Corradini, P., 1960. Nuevo Cimetno Suppl. 15 (10), 40. Nishida, K., Okada, K., Asakawa, H., Matsuba, G., Kazuki, I., Kanaya, T., et al., 2012. Polym. J. 44, 95. Padden, F.J., Keith, H.D., 1959. J. Appl. Phys. 30 (10), 1479. Pasquini, N. (Ed.), 2005. Polypropylene Handbook. Carl Hanser Verlag, Munich. Passaglia, E., and Kevorkian, H.K., 1963. J. Appl. Phys. 34, 90. Roberts, J.S., 1982. U.S. Patent 4,347, 206. Samuels, R.J., 1970. J. Macromol. Sci., Phys. B4, 701. Samuels, R.J., 1975a. J. Polym. Sci., Polym. Phys. Ed. 13, 1417. Samuels, R.J., 1975b. Appl. Polym. Symp. 27, 205. Samuels, R.J., J. Polym. Sci., 1979. Polym. Phys. Ed. 17, 535. Sauer, J.A., Pae, K.D., 1968. J. Appl. Phys. 39, 4959. Schaefer, D., Spiess, H.W., Suter, U.W., Fleming, W.W., 1990. Macromolecules 23, 3431. Schick, C., 2018. Private communication to Menczel, J. Schick, C., Androsch, R., Schmelzer, J.W.P., 2017. J. Phys. Condens. Matter. 29, 453002 (35pp). Sheehan, W.C., Cole, T.B., 1964. J. Appl. Polym. Sci. 8, 2359. Silvestre, C., Cimmino, S., Duraccio, D., Schick, C., 2007. Macromol. Rapid Commun. 28, 875. Steinmann, W., Walter, S., Beckers, M., Seide, G., Gries, T., 2013. Thermal analysis of phase transition in polymeric fibers, Chapter 12. In: Elkordy, A.A. (Ed.), Applications of Calorimetry in a Wide Context, Differential Scanning Claorimetry, Isothermal Titration Calorimetry and Microcalorimetry. IntechOpen, 2013. Tashiro, K., Naki, Y., Kobayashi, M., Tadokoro, H., 1980. Macromolecules 13, 137. Thierry, A., Fillon, B., Straupe´, C., Lotz, B., Wittmann, J., 1992. Prog. Colloid Polym. Sci. 87, 28. Todoki, M., Kawaguchi, T., 1977a. J. Polym. Sci., Polym. Phys. Ed. 15, 1067. Todoki, M., Kawaguchi, T., 1977b. J. Polym. Sci., Polym. Phys. Ed. 15, 1507. Turner-Jones, A., 1971. Polymer. (Guildf). 12, 487. Varga, J., 1995. Crystallization, melting and supermolecular structure of isotactic: polypropylene. In: Karger-Kocsis, J. (Ed.), Polypropylene: Structure, Blends and Composites, vol. 1. Chapmann & Hall, London, p. 56. Weeks, N.E., Porter, R.S., 1975. J. Polym. Sci., Polym. Phys. Ed. 13, 1177. Wunderlich, B., 1976. Macromolecular Physics, Volume 2, Crystal Nucleation, Growth, Annealing. Academic Press, New York, San Francisco, London. Wunderlich, B., Macromolecular Physics, Volume 3, Crystal Melting, 1980, Academic Press, New York, 1980.

222

Thermal Analysis of Textiles and Fibers

Further reading Andreassen, E., Grostad, K., Myhre, O.J., Braathen, M.D., Hinrichsen, E.L., Syre, A.M.V., et al., 1995. J. Appl. Polym. Sci. 57 (9), 1075. De Santis, F., Adamovsky, S., Titomanlio, G., Schick, C., 2006. Macromolecules 39 (7), 2562. Gedde, U.W., 1999. Polymer Physics. Kluwer Academic Publisher, Dordrecht. Horva´th, Z., Menyha´rd, A., Doshev, P., Gahleitner, M., Tranninger, C., Kheirandish, S., et al., 2013. J. Appl. Polym. Sci. 130 (5), 3365. ´ .O., Horva´th, Z., et al., 2015. Menyha´rd, A., Suba, P., La´szlo´, Z., Fekete, H.M., Mester, A Express Polym. Lett. 9 (3), 308. Monasse, B., Haudin, J.M., 1985. Colloid Polym. Sci. 263 (10), 822. Monks, A.W., White, H.M., Bassett, D.C., 1996. Polymer. (Guildf). 37 (26), 5933. Nishida, K., Okada, K., Asakawa, H., Matsuba, G., Ito, K., Kanaya, T., and Kaji, K., 2012. Polym. J. 44, 95. Riley, J.L., 1956. In: Schildknecht, C.E. (Ed.), Spinning and Drawing Fibers. Wiley Interscience, New York, Chapter XVIII. Romankiewicz, A., Sterzynski, T., 2002. The lamellar distribution in isotactic polypropylene modified by nucleation and processing. Macromol. Symp. 180, 241 256. Samuels, R.J., 1973. J. Macromol. Sci. B8 (1 21), 41. Samuels, R.J., 1979. J. Polym. Sci., Polym. Phys. Ed. 15, 535. Schick, C., Wurm, A., Mohamed, A., 2002. Thermochim. Acta 392, 303. Turi, E.A., 1981. Thermal Characterization of Polymeric Materials. Academic Press Inc, New York. Varga, J., Menczel, J., Solti, A., 1981. J. Therm. Anal. 20, 23. Varga, J., Fujiwara, Y., Ille, A., 1990. Period. Polyt. Chem. Eng. 34 (4), 255. Varga, J., 2002. J. Macromol. Sci. Phys. B41 (4 6), 1121. Wurm, A., Schick, C., 2003. Colloid Polym. Sci. 281 (2), 113 122 (2003).

Thermal analysis of aliphatic nylon fibers

13

Lawrence Judovits Arkema, King of Prussia, PA, United States

Abstract Nylon is the common name for polyamides. The melting points for nylons made from amino acids generally increase as the amide content increases although exceptions exist. For example, odd-numbered nylons have a higher melting point in its sequential trend. Nylons make good fibers in part due to the intermolecular attractions associated with hydrogen bonding between the polyamide linkages. The hydrogen bonding also provides readily available nucleation sites. The glass transition temperature for nylons ranges from 40
C to 60
C. Nylon fibers are known to undergo reorganization in a differential scanning calorimetry experiment where the original melting peak increases upon heating. Stress release of nylon fibers can also be followed using thermomechanical analysis by monitoring the shrinkage and rate of shrinkage. Nylon-6 and 6,6 resins account for most of the nylon fiber production. Nylon fibers are ubiquitous and found in many applications, such as carpets, textiles, threads, fishing lines, and tire cords. Top producers of nylon fibers are INVISTA, Formosa Chemicals and Fibre Corporation, Shaw Industries, Fujian Jinjiang Technology, and Zhejiang-based Yiwu Huading Nylon.

13.1

Nylon fiber production and basics

Nylon is a common name for the family of polyamide resins with nylon-6 and nylon6,6 being the largest in terms of consumption, which accounts for most of the fiber production (Sriram et al., 2013). Global nylon fiber capacity declined in the last decade, associated with the Great Recession of the 2000s, though the world demand for nylon fiber in 2013 was 3.9 million metric tons with Asia continuing to lead in spite of the production being much less than the world capacity of 7.2 million metric tons (Davis et al., 2014). The top producers of nylon fiber in 2014 were INVISTA, Formosa Chemicals and Fibre Corporation, Shaw Industries, Fujian Jinjiang Technology, and Zhejiang-based Yiwu Huading Nylon. However, they only account for 18.8% of the total global production indicating many small producers as illustrated later (Fig. 13.1). Most of the production of nylon fiber occurs in Asia and in 2013 accounted for two-thirds of the global output. The world consumption of nylon fibers can be divided into the following five areas according to the Chemical Economics Handbook on Nylon Fibers (Davis et al., 2014): 1. textile filament yarn (47.4% of global consumption); 2. industrial filament yarn (30.9%); Thermal Analysis of Textiles and Fibers. DOI: https://doi.org/10.1016/B978-0-08-100572-9.00013-6 © 2020 Elsevier Ltd. All rights reserved.

224

Thermal Analysis of Textiles and Fibers

INVISTA, 424,000 metric tons FCFC, 262,000 metric tons Shaw Industries, 250,000 metric tons JJ, 206,000 metric tons Huading, 206,000 metric tons Rest of the world, 5,857,000 metric tons

Figure 13.1 Top producers of nylon fibers for 2014. Source: Data from Davis, S., Grayson, R., Yalcinkaya, M., Wu, M., Yoneyama, M., 2014. Chemical Economics Handbook, Nylon Fibers, IHS Chemicals. 3. bulk continuous fiber (BCF) (17.4%); 4. staple (4.2%); and 5. other fiber types, such as monofilament, which only accounted for 0.1% of the global consumption.

Nylon fibers are used in a variety of applications such as carpets and rugs, apparel and hosiery, tire cord and fabric, belting and hose, textiles and wovens, rope and cordage, sewing thread, and fishing line. Yarns composed of continuous filaments may be used for textiles or other industrial uses. Either BCF or staple (fiber cut to a set length) can be used to produce carpets. Tire cord is produced from high-tenacity continuous-filament fiber yarn (Tyre Cord, 2015). Nylon monofilament is used in translucent fishing lines and as nonabsorbable sutures (ABC-ofFishing, 2015; Dolphin Sutures, 2015). Of the five categories, by far, textile yarn is the largest making up almost 50% of the consumption of nylon fiber in 2013. Nylon fiber is defined as “A manufactured fiber in which the fiber-forming substance is a long-chain synthetic polyamide in which less than 85 percent of the amide linkages are attached directly to two aromatic rings” by the US Federal Trade Commission (Electronic Code of Federal Regulations, 2014). This definition can be considered broad since the definition is based on the type of bond, not a particular monomer. Although as mentioned, nylon-6 and nylon-6,6 account for most of the nylon fiber production, approximately 95%98%; other nylons are also used for fibers, which include nylon-4,6; 6,10; 6,12; 11; 12; and aramids (Davis et al., 2014). Both nylon-6 and nylon-6,6 were invented in the late (c.) 1930s, with nylon6,6 being discovered first by Wallace Carrothers (1938), who was working for DuPont, and then nylon-6 followed by Schlack (1941) at IG Farben. Nomenclature for nylon indicates whether an amino acid, denoted as a single number, or a diamine and diacid, denoted as a double number, separated by a comma, are used in the polymerization to make the homopolymer (Microlab, 2015). Use of a numbered notation separated by slashes would indicate a copolymer, such as nylon-6/6,6, which would be composed of an amino acid, a diacid, and a diamine. The number designates the number of carbon atoms in each

Thermal analysis of aliphatic nylon fibers

225

Figure 13.2 An example of hydrogen bonding in nylons.

monomer unit. Source-based nomenclature, for example, poly(ε-caprolactam) instead of nylon-6, will not be used in this section. The fact that nylons make good fibers is due, in part, to the intermolecular forces resulting from the hydrogen bonding between the polyamide linkages, that is, the nitrogen-bonded hydrogen of one polymer chain can bond with the carbonyl oxygen of an adjacent chain as depicted in Fig. 13.2. The hydrogen bonding not only increases the crystal strength but also allows the chains to align when formed into a fiber. Nylon crystallinity for unoriented samples has been reported from 20% to 30% (Greco and Nicolais, 1976) (although with other thermal treatments, a wider range in crystallinities would be expected), so drawing will additionally increase the crystallinity along with orienting the amorphous regions, and hence the fiber strength will also increase.

13.2

Thermal analysis basics of nylon and its fibers

As with any polymer, commercial nylon resins and fibers typically melt well below their equilibrium melting point, and this depression is generally associated with their crystal size and perfection (Wunderlich, 1980) although trends can be noted between the different nylon grades. Based on data on nylon melting points, measured by several techniques, the following was noted in general (Starkweather, 1995; Miller, 1989): 1. For separate even- and odd-series nylons, higher amide content (lower nylon number) has a higher melting point (e.g., see Table 13.1 for equilibrium melting points of nylon-6 and nylon-6,6). 2. Odd-numbered nylons have a higher melting point compared to their sequential evennumbered nylons. 3. Double-unit eveneven-series nylons have higher melting points compared to their single-unit analogs, such as nylon-6,6, which has a higher melting point than nylon-6. 4. Extrapolation of the various series to the zero amide condition gives a melting point lower than that of polyethylene, which Starkweather et al. (1984) attribute to the nylon methylene groups not achieving optimum crystal packing (in the space occupied in a crystal lattice).

226

Thermal Analysis of Textiles and Fibers

Table 13.1 Equilibrium melting data and glass transition temperature (10
C/min cooling and heating) (Xenopoulus, 1990). Material

Tm0 (
C)

ΔHf0 (kJ/mol)

ΔHf0 (J/g)

Tg (
C)

ΔCp,100%a (J/
C mol)

ΔCp,100% (J/
C g)

Nylon-6 Nylon-6,6 Nylon-11 Nylon-12

260 301 220 227

25.0 57.8 44.7 48.4

221 255 244 245

40 50 43 41

54 112 69 74

0.48 0.49 0.38 0.38

a

ΔCp,100% is the heat capacity jump at the glass transition for the totally amorphous sample.

The higher melting point of nylons from that of polyethylene has been attributed to their low entropies of fusion, which is due to hydrogen bonding that still occurs in the melt (Dole and Wunderlich, 1959; Starkweather et al., 1984). Conversely, the hydrogen bonding provides readily available nucleation sites for crystallization. Polymer “crystalline memory” refers to the effect of the thermal and mechanical history prior to the solid state, which influences the behavior of the polymer even after it is cooled from the melt (Jaffe et al., 1997). This would mean that the crystallization from the melt would also be a function of the prior annealing temperature or mechanical stresses imposed on the polymer if those stresses were not relieved. The memory effect in nylon can be attributed to the persistence of crystal nucleating species, which can be deactivated by annealing for an extended time at temperatures in the melt above the equilibrium melting temperature. These species are thought of as molecular clusters, which maintain their crystallographic identity that is usually caused by, for example, partial melting. If caused by mechanical deformation, this is further referred to as an orientation memory effect (Supaphol and Lin, 2001). For nylon-6, annealing at 280
C for, at least, 90 minutes has been suggested to eliminate memory effects (Avramova and Fakirov, 1986). Generally in DSC studies, glass transition temperatures were found to be ranging from 40
C to 60
C (Greco and Nicolais, 1976; Xenopoulus, 1990). Vetted equilibrium melting points and heat of fusion as well as glass transition temperature data for nylon-6, nylon-11, nylon-12, and nylon-6,6 have been analyzed and given in Table 13.1 (Xenopoulus, 1990). Determination of equilibrium melting points for polyamides that were obtained on annealed cross-linked drawn fibers was excluded on general concerns of deviations caused by cross-linking on crystallite size when extrapolating to the equilibrium condition using the GibbsThomson equation (Xenopoulus, 1990). The GibbsThomson equation relates the size of a crystal (for a polymer, its lamellar thickness) to its melting point so that the equilibrium melting point can be determined by extrapolation. Complicating any thermal analysis of nylons is cold crystallization, the Brill transition, and reorganization, which have been shown to occur in these systems. Cold crystallization and crystallization on heating have been noted in a wide variety of nylons when heating above the glass transition (see Fig. 13.3) (Illers, 1977); however, it is rarely seen in drawn fibers. One would assume that the additional

Thermal analysis of aliphatic nylon fibers

227

Figure 13.3 DSC scans of quenched polyamides heated at 20
C/min: (A) nylon-6, (B) nylon-7, (C) nylon-8, (D) nylon-9, (E) nylon-11, (F) nylon-12, and (G) nylon-15. Source: Reproduced from Illers, K.H., 1977. Polymer 18, 551 (Fig. 1) with permission from Elsevier.

crystallization that occurs during drawing results in lowering the amorphous content thereby reducing any cold crystallization. The Brill transition is a reversible polymorphic phase transition, which may be observed just below the nylon melting. The Brill transition can be determined through wide-angle X-ray diffraction (WAXD) when the interchain distance, within a hydrogen-bonded sheet, and the intersheet distance, between hydrogen-bonded sheets, merge to a single reflection (Vinken, 2008). The Brill transition is difficult to note by DSC although it can be seen as a distinct endotherm in some solutioncrystallized specimens (Starkweather and Jones, 1981). For nylon-11 a Brill transition can occur not only as the temperature is increased, but also after being elongated greater than 20% (Dang and Tence-Girault, 2008). This additional deformation before fracture has been attributed to an energy dissipation mechanism resulting in better impact resistance (Dang and Tence-Girault, 2008). Reorganization (Wunderlich, 1980) is where the initial metastable crystals will further perfect themselves upon heating. This may occur through either an annealing or a melting/recrystallization process and usually results in multiple peaks. Reorganizational behavior in nylon appears to have been initially suggested by Wilhoit and Dole (1953) where they theorized that this occurred for drawn filaments of nylon-6,6 that melted at a lower temperature than the corresponding undrawn filaments. Multiple peak melting behavior is thought to be first experimentally reported by White in 1955 where a doublet melting peak was noticed in a differential thermal analysis (DTA) trace for nylon-6,6. Reorganization in nylon-6 was

228

Thermal Analysis of Textiles and Fibers

dH/dt

228 225 Control

222

1 mcal/sec

219 216

144 164

Control (unirradiated)

212

186 207

180

190 200 210 220 Temperature (°C)

230

Figure 13.4 Inhibited reorganization in nylon-6 by cross-linking. Source: Reproduced from Todoki, M., Kawaguchi, T., 1977b. J. Polym. Sci. Polym. Phys. Ed., 15, 1067 (Fig. 1) with permission from Wiley.

demonstrated by Todoki and Kawaguchi (1977b) on studies of drawn yarn samples by DSC, which were then annealed and cross-linked. Cross-linking was induced by infusing the nylon with acetylene and then irradiating with γ radiation. The acetylene acts as a cross-linker, and the reaction takes place in the nylon amorphous phase. An unannealed, unirradiated sample, which acted as a control, showed a doublet melting peak at high temperatures, while the irradiated samples showed a single melting peak lower in temperature than the control but higher than their annealing temperature, its reorganization inhibited by the cross-linking (Fig. 13.4). Similar results were obtained by Arakawa (1968) on drawn nylon-6 that had been subjected to methoxymethylation of amide groups in the amorphous region. By depriving amorphous segments of their mobility, reorganization was inhibited, resulting in zero entropy production melting (Wunderlich, 1980). Zero entropy production melting is where the metastable crystal melts directly into a melt of the same metastability below the equilibrium melting point. Another method to observe the original morphology or differentiate between different thermal histories is to heat fast and, in doing so, retard or inhibit reorganization. Such a methodology has been demonstrated for forensic applications in fingerprinting the source/manufacturer of the fiber (PerkinElmer, 2015). Todoki and Kawaguchi (1977a) also investigated constrained nylon fibers and found a single melting peak at temperatures higher than that of the unconstrained (or free-to-shrink) melting. The constrained fibers were held at constant length during heating by winding around a metal plate with the ends tied. The unconstrained fiber, which is free to shrink, showed multiple melting peaks indicative of reorganization. To compare each of the processing conditions, Table 13.2 gives the lastobserved peak temperature upon heating. Observed in Table 13.2 is an increased

Thermal analysis of aliphatic nylon fibers

229

Table 13.2 Comparison between unconstrained, constrained, and cross-linked nylon-6 fibers (Todoki and Kawaguchi, 1977a). Melting peak temperature (
C)

Drawing conditions Temperature (
C)

Ratio

Cross-linked

Unconstrained

Constrained

Undrawn (as-spun) 55

 2.5 3.0 4.0 5.0 2.5 3.0 4.0 5.0 2.5 3.0 4.0 4.6

151 156 158 160 167 173 172 172 177 197 196 199 195

221.0 220.5 220.8 221.3 222.2 220.0 220.8 221.2 222.1 220.5 220.7 221.7 221.8

222.0 223.5 225.2 229.1 231.2 223.9 225.5 228.1 232.4 223.5 225.4 228.2 230.3

120

170

Only the last-observed melting peak upon heating is reported.

melting peak temperature with the draw ratio for the constrained fiber. This increased melting peak is probably due to superheating since the melt is also held oriented (due to being constrained). The constrained fiber gave the highest melting peak temperature when compared to the unconstrained or radiation cross-linked samples made under the same processing conditions. Semicrystallinity is used to define the structure of polymers where two phases are present, crystalline and amorphous. Semicrystalline polymers are believed to crystallize into lamella stacks separated by an amorphous phase; however, the amorphous phase itself has been found to be complicated, for example, two different types of segmental motion have been observed in the amorphous phase of nylon-6,6 by solid-state deuterium NMR (nuclear magnetic resonance) (Miura et al., 1990). Murthy (2006) also noted that a simple semicrystalline model does not fully describe the amorphous morphology of nylons and its fibers nor does it fully explain its mechanical properties especially after drawing. Since one can find varying degrees of order in the crystalline phase, it can therefore be argued that the amorphous phase may also have varying degrees of disorder (Murthy et al., 1995). Murthy (2006) suggests that the packing density of the amorphous chains is not uniform and different components become differently oriented upon drawing. Also, complicating the amorphous structure is the existence of a rigid amorphous fraction (RAF) similar to what has been found for polyethylene terephthalate fibers (Fu et al., 1994). One therefore needs to differentiate between oriented amorphous morphology, which can exist independent of any crystallinity, and RAF, which exists at the interphase between the bulk amorphous and its crystalline phase. Oriented amorphous morphology can relieve itself when the material is heated above the glass transition, while RAF can remain up to the melting point.

230

Thermal Analysis of Textiles and Fibers

In composites’ transcrystallinity, a morphology, different from the matrix, can occur at the interface or form an interphase between the filler and the resin matrix (Clark, 1996). Transcrystallinity was originally observed in nylon-6,6 by Barriault and Gronholz (1955) through optical microscopy and was found to form under nearly all conditions except severe quenching. Transcrystallinity, being semicrystalline, has an associated amorphous phase with a possible orientation, which would then be separate from the matrix. Thermomechanical analysis (TMA) is typically used to measure the coefficient of linear thermal expansion; however, for fibers, it is important to study shrinkage though elongation may also be of interest. The fiber response depends on the inherent tendency of its internal stress to relieve itself when permitted by a loss of the resisting morphological structure (e.g., glass or crystal). TMA measurements are generally done under some load where a low force is used to study shrinkage and thermal expansion while a high force is used to study elongation. Nylon fibers, when analyzed axially, typically exhibit a two-stage shrinkage with temperature, which has been attributed to the relaxation of the oriented amorphous chains at low temperature, while at higher temperature crystal reorganization (Murthy, 2006). Dynamic mechanical analysis (DMA) of common aliphatic nylons generally shows three loss peaks below the melting transition. The glass transition or α relaxation is found at the highest temperature, a smaller peak, the β transition, at about 240
C, which is greatly affected by water content, while the third transition, the γ relaxation, has a peak around 2120
C and is probably due to the localized motion of the CH2 groups (Ong et al., 1986; Serpe and Chaupart, 1996; Starkweather, 1995). The α peak is related to the motion associated with the breaking of the hydrogen bonding to the CONH groups (Serpe and Chaupart, 1996). Starkweather (1995) gives loss peaks for aliphatic nylons ranging from 53
C to 80
C with the lowest given for nylon-6 and the highest for nylon-6,6. These transitions can be seen in Fig. 13.5 with the glass transition peak higher and narrower for the higher methylene content. The glass transition by DMA is usually higher than that from DSC due to frequency effects. All the transitions shift to lower transitions as the amount of moisture, which acts as a plasticizer, is increased although one could assume that other solvents, such as some alcohols, which interrupt the hydrogen bonding, will also affect these transitions (Starkweather, 1995). Thermal decomposition of aliphatic nylons has been noted to proceed by a variety of competing mechanisms resulting in chain scission probably at the NHCH2 bond (Levchik et al., 1999). The suggested mechanisms include hydrolysis, homolytic scission, and intramolecular hydrogen transfer. Thermo-oxidative decomposition appears to start either at the methylene group vicinal to the amide unit or at the β position to the carboxyl methylene group. Molten resin exposed to air can crosslink and gel, causing problems with fiber spinning (Levchik et al., 1999). Photolytic decomposition from environmental exposure is also known although light stabilizer packages can be used (Pagilagan, 1995; Rilsan, 2015). In general, polylactams can decompose back to their cyclic monomers, while the products formed from diaminesdiacids are more complicated. Thermogravimetric analysis (TGA) is commonly used to study nylons and common examples are given in Figs. 13.6 and 13.7 although differences between nylon sources may affect the particular resin stability as many nylons are heat stabilized (Rilsan, 2015).

Thermal analysis of aliphatic nylon fibers

231

–0.70

–0.88 Log –1.07 tan δ –1.26

–1.44

–1.63

–1.81 –2.00 –120

–64

–8 48 104 Temperature (ºC)

160

Figure 13.5 DMA at 110 Hz for ’ nylon-6; K nylon-11; ▲ nylon-12; V nylon-6,6; ▼ nylon-6,12; nylon-6,10; and nylon-6,6/6. Source: Reproduced from Ong, E.S., Kim, Y. Williams, H.L., 1986. J. Appl. Polym. Sci. 31, 367 (Fig. 1) with permission from Wiley.

Figure 13.6 TGA using air as the sample purge at 20
C/min performed on a TA Instruments Q5000IR. The nylon-6 and nylon-6,6 pellets were from Polysciences, while the nylon-12 resin was from Arkema Inc. (Judovits, to be published).

232

Thermal Analysis of Textiles and Fibers

Figure 13.7 (A)TGA using nitrogen as the sample purge at 20
C/min performed on a TA Instruments Q5000IR. The nylon-6 and nylon-6,6 pellets were from Polysciences, while the nylon-12 resin was from Arkema, Inc. (Judovits, to be published). (B). TGA of nylon-11 (solid line) and nylon-12 (dashed line) under nitrogen at 5
C/min performed on a TA Instruments 2950. The nylon-11 sample used was obtained from Arkema, while the nylon-12 sample was from Sigma-Aldrich. Source: Reproduced from Telen, L., Puyvelde, P.V., Goderis, B., 2016. Macromolecules 49, 876 (Fig. 1) with permission from ACS Publications.

13.3

Nylon-6

Nylon-6 has two commonly found crystalline phases denoted α and γ, both monoclinic, with the α phase being the more thermodynamically favored while the γ phase the more kinetically favored (Murthy, 2006; Wunderlich, 1973). Rapid

Thermal analysis of aliphatic nylon fibers

233

Table 13.3 Physical properties of nylon-6 polymorphs. Nylon

Form

Density (g/cm3)

Young’s modulus (GPa)

References

Nylon-6

α, Monoclinic γ, Monoclinic Amorphous

1.23 1.17 1.08

295 135 

Arimoto et al. (1965); Li and Goddard (2002); Wunderlich (1973)

crystallization from the melt favors the γ form, and slow crystallization favors the α form. The α-form crystal is composed of extended zigzag chains with antiparallel hydrogen-bonded sheets, while the γ-form crystal is composed of pleated chains with parallel hydrogen bonding (Gianchandani et al., 1982). The two crystalline forms are known to coexist in nylon-6 fibers in various percentages depending on processing conditions. Since the Young’s modulus of the α-form crystal is higher than that of the γ-form crystal, the different mixing of the α and γ crystalline forms impart different mechanical properties (see Table 13.3). In this regard, nylon-6 appears to be a special case. For aliphatic nylons with less than six carbons, the α form is the more stable, while for those with more than six carbons, the γ form is more stable. This results for nylon-6 the likelihood of two coexisting forms (Dasgupta et al., 1996). Nylon-6 fibers are typically produced by melt spinning followed by drawing. Drawing is affected by the draw ratio, temperature, take-up speed, the number of drawing stages, and heat-set with nylon-6 typically reaching a crystallinity of 50% (Murthy et al., 1995). Since nylon-6 has a low-enough glass transition, it can be cold-drawn in addition to being processed by hot drawing. Each fiber process uniquely affects the amorphous and crystalline structure that is formed. Nylon-6 fiber is typically cold-drawn for apparel and carpet, whereas for tire cords and car seat belts, where high fiber strength is needed, it is subjected to hot multistage drawing (Rahbar and Mojtahedi, 2011). As pointed out earlier in this chapter, nylon-6 undergoes reorganization when heated but when cross-linked, the original crystal morphology could be observed (Todoki and Kawaguchi, 1977b). Similar work by Arakawa (1968) and Arakawa et al. (1969a,b) found that methoxymethylation of amide groups also cross-linked the amorphous regions of drawn nylon-6 fibers. By hindering or depriving amorphous segments, the ability to crystallize retarded or prevented reorganization. Drawn, 3.2 times ( 3 ) the original length, nylon-6 filaments were annealed at 170
C for 1 hour and then methoxymethylated at progressive times. A decrease in melting temperature, from the as-annealed sample, from 220
C to 195
C was observed after 30 hours of reaction after which there was no further change with longer reaction times. The shift from a high melting point to a low melting point was repeated for annealing temperatures between 150
C and 215
C. The reaction time required to obtain a stable peak was dependent on the annealing temperature. In Table 13.4 the methoxymethylation method was used to study the effect of increasing annealing temperature followed by cross-linking for a reaction time that

234

Thermal Analysis of Textiles and Fibers

Table 13.4 Nylon-6 filaments after annealing for 1 h and methoxymethylation (Arakawa et al., 1969b). Annealing temperature (
C)

Long period ˚) (A

˚) Thickness (A using density

˚) Thickness (A using ΔHf

Melting point (
C)

150 155 160 165 170 175 180 185 190 195 200 205 210 215

80.1 80.2 81.5 81.5 84.7 86.8 88.3 92.1 94.5 98.1 102 104 111 117.5

24.7 25.2 26.8 26.9 29.6 31.6 33.0 34.6 39.4 43.5 48.0 50.7 58.8 67.5

22.8 23.5 24.4 25.1 26.6 28.1 29.3 32.4 35.1 39.2 43.8 46.8 52.1 57.8

193.7 194.4 195.0 194.9 195.8 197.5 200.1 202.9 205.4 210.1 214.6 215.4 217.7 221.1

gave a stable peak. The thickness of the crystallites was calculated from the long period (thickness of the crystallite and its associated interface) adjusting for the amount lost to amorphous material, between the lamellar stacks, by using crystallinity values from either density or heat-of-fusion measurements. A correlation of increasing lamella thickness to melting point can be noticed although the extrapolated equilibrium melting point is higher than expected, suggesting that high amounts of cross-linking may result in superheating. Koyama et al. (1986) investigated the effect of molecular weight for the highspeed spinning of nylon-6. In general, the crystallinity and the γ fraction were both found to increase with increasing take-up velocity; however, the crystallinity at high take-up velocity decreased with increasing molecular weight, while the converse was noted at low take-up velocity. This was attributed to competing effects on the crystallization rate versus the final attainable crystallinity. The highmolecular-weight samples, spun at high speeds, also contained a larger proportion of α-phase crystals than did low-molecular-weight samples, which most likely resulted from a phase transformation from the γ phase at elevated temperature on the running thread line. Nylon-6 fibers drawn under high pressure were investigated by Miyata and Balakov (1981). Their fiber drawing was done at 80
C using a draw ratio up to 3.75 times more and pressures from atmospheric to 2000 kg/cm2. They found a conversion from the γ to the α form with increasing pressure. Dennis and Buchanan (1987) used two different thermomechanical methods to study shrinkage and shrinkage forces on nylon-6 carpet yarns by classical TMA and thermal stress analysis (TSA), respectively. Classical TMA experiments can provide information regarding the expected shrinkage during the manufacturing process or during subsequent use, while TSA can monitor the development of thermal stresses

Thermal analysis of aliphatic nylon fibers

235

or shrinkage forces with temperature. Their results were consistent with two mechanisms similar to what has been described by Murthy (2006), the first due to entropic considerations, such as oriented amorphous stress release, and the second associated with crystallization or reorganization. Khanna (1990) studied oriented nylon-6 annealed and conditioned by three different processes to heat-set the fiber: the Suessen method (which uses dry heat), in an autoclave (which used saturated steam in a batch process), and the Superba method (in a sense a continuous autoclave process where saturated steam is used). He found that if the yarn was subjected to a heat treatment, three instead of two endotherms appear. Typically, for moisture-containing nylon-6 yarn, two endotherms can be observed with the lowest representing the loss of moisture, while the second is due to melting. Heat-setting results in a middle endothermic peak (MEP) due to annealing during the heat-set process. MEP was found to increase linearly with the heat-setting temperature. Khanna also theorized that the MEP represented melting of the crystallites that formed in the amorphous regions of the fiber microstructure during the heat-setting process. By TMA, Khanna (1990) found that the heat-set yarn showed less shrinkage when compared with an unconditioned sample, and a linear correlation exists between the heat-setting temperature and the onset temperature of shrinkage below the main melting. He attributed this, again, to originate from the formation of small crystallites. Accelerated shrinkage of the nylon-6 yarns occurred at a temperature above 180
C to about 220
C where main crystal melting began. Kunugi et al. (1982a,b,c) used a zone-drawing and a zone-annealing method followed later by a multistep procedure to prepare a high-modulus, high-strength nylon-6 fiber. The multistep method resulted in a tensile storage modulus 1.5 times greater than that of the single-step zone-annealing method, which was 2.24 times greater than that of the high-tenacity fibers available commercially. DMA indicated that amorphous chain motion became more hindered for the zone-annealed fiber causing an increase in the α-peak temperature compared to other preparations as noted in Table 13.5. As noted in Fig. 13.8, the annealing after drawing increased the modulus more than just drawing. They attributed the high modulus to the large Table 13.5 Dynamic mechanical analysis of α-relaxation peak temperatures for different thermal treatments (Kunugi et al., 1982b). Sample Fiber annealed under release Fiber annealed at constant length Fiber annealed under tension Fiber zone-annealed

Ev, α peak (
C)

tan δ, α peak (
C)

Young’s modulus ( 3 10210 dyne/cm2)

75

75

3.71

78

83

4.12

50

88

5.73

96

95

9.91

236

Thermal Analysis of Textiles and Fibers

Figure 13.8 DMA storage modulus for fiber annealed under release ▲, fiber annealed under constant length Δ, fibers annealed under tension K, and zone-annealed fiber x. Source: Reproduced from Kunugi, T., Akiyama, I., Hashimoto, M., Matsuzaki, K., 1982c. Polymer 23, 1983 (Fig. 1) with permission from Elsevier.

number of tie molecules connecting the crystallites as well as to the high orientation of the amorphous chains. Based on their DMA studies of nylon-6, Rong and Williams (1985) found that water acts as a plasticizer above 210
C but below as an antiplasticizer as seen in the storage modulus. They attributed the antiplasticization to water acting as a strong bonding agent with the nylon dipoles. The glass transition peak or α relaxation of the drawn fiber splits into two with the larger portion, denoted by α a decreasing in temperatures with increasing water content, whereas the smaller higher temperature peak, now denoted by αa, remains essentially fixed at about 90
C. This can be seen in Table 13.6. Two other relaxations were noted, β and γ, which are also plasticized with moisture. The β relaxation ranged from 227
C to 250
C, while the γ relaxation ranged from 2105
C to 2114
C with the lower end of the range favored by higher moisture. La Mantia et al. (Acierno et al., 1977; La Mantia et al., 1990) have studied various substances added to nylon-6 such as salts and liquid crystal copolymers. The salt-containing nylon-6 fiber was investigated by DMA, and the glass transition was determined as a function of draw ratio (La Mantia et al., 1982). The addition of LiCl was found to reduce the crystallization rate of nylon-6 while increasing the melt viscosity due to a cross-linking effect resulting from polymersalt interaction. This is different from the addition of a liquid crystal copolymer (Vectra B950), which decreased the melt viscosity but increased the elastic modulus (La Mantia et al., 1990). Structural determinations of the nylon-6 LiCl systems, obtained through density, birefringence, and WAXD measurements, found that it was possible to achieve a high amorphous orientation. The glass transition temperature was found to increase with the addition of the salt and is explained due to pseudo-crosslinking between lithium ions and the oxygen on the polyamide carbonyl group (La Mantia, and Acierno, 1981). Another approach to overcome the quasi-cross-linking, from the intermolecular hydrogen bonding, was that of reversible plasticization

Thermal analysis of aliphatic nylon fibers

237

Table 13.6 Effect of relative humidity (RH) on the α peak (Rong and Williams, 1985). α a

αa RH (%)

tan δ




0 8 33 75 97 Swollen

0.129 0.133 0.137 0.133 0.137 0.135

93 89 90 91 100 99

C

tan δ




No peak 0.124 0.141 0.168 0.171 0.176

No peak 45 32 20 18 6

C

using liquid ammonia that is volatized upon exiting the die during extrusion (Zachariades and Porter, 1979). About 18 mass% of ammonia was determined by TGA after presoaking under pressure but before extrusion. Liu et al. (2007) prepared nylon-6 nanofibers with diameters around 200 nm through electrospinning. They studied the as-electrospun fiber and compared it to a solution precipitant and a polyimide-coated nylon-6. The solution precipitant was prepared by precipitating the polymer from the electrospinning solution of 8 mass% nylon-6 in 1,1,1,3,3,3-hexafluoro-2-propanol using diethyl ether as the nonsolvent followed by drying to a constant mass. In Fig. 13.9 the solution precipitated from the material, which was used in electrospinning, showed reorganization by having the highest melting peak, while the confined fiber showed a low-temperature solder indicative of poorly formed crystals that could not reorganize. The electrospun fiber demonstrated even more imperfection so that cold crystallization could be observed. As with melt-spun nylon-6 fibers, electrospun nylon-6 fibers exhibited predominantly the γ crystalline form with the chain axis preferentially oriented parallel to the fiber axis. Upon annealing above 150
C, γ-form crystals converted into the more thermodynamically stable α-form crystal. For confined polyimide-coated fibers, the γ-form was inhibited from transforming into the α-form crystals until 210
C, only 10
C below its melting peak point at 220
C. The Brill transition for the polyimide-confined nanofibers was observed at 180
C190
C through WAXD, which was at least 20
C higher than that in unconfined nylon-6 at approximately 160
C and it also attributed to the confinement effect. Gupta et al. (2009) obtained electrospun porous nylon-6 fibers from Lewis acidbase complexation with gallium trichloride (GaCl3) followed by washing out the GaCl3. The unwashed complex metal fibers did not show any melting up to 230
C, suggesting an amorphous structure, while the GaCl3 water-washed fiber demonstrated a typical nylon-6 melting. GaCl3, being a Lewis acid, interacts with the Lewis base carbonyl sites, which may prevent the chains from crystallizing. Wang et al. (2012) produced electrospun nylon-6 fibers from formic acid solutions at different concentrations. Two distinct melting peaks were observed and referred to as the lower melting temperature (LMT) and higher melting temperature (HMT) phases with the HMT peak melting above the equilibrium melting. As noted in Fig. 13.10, variable heating rates do not show an interconversion between the

238

Thermal Analysis of Textiles and Fibers

Figure 13.9 DSC of electrospun fiber (solid line), its solution precipitated as a reference (dashed line), and polyimide-coated nylon-6 nanofiber (dash-dotted line) heated at 10
C/min. Source: Reproduced from Liu, Y., Cui, L., Guan, F., Gao, Y., Hedin, N.E., Zhu, L., et al., 2007. Macromolecules 40 (17) 6283 (Fig. 3) with permission from ACS Publications.

Figure 13.10 DSC of electrospun nylon-6 fibers at different heating rates. The fibers were obtained from a 22-mass% nylon-6 formic acid solution. (A) represents the DSC scans with different heating rates; (B) is the change in melting temperature with heating rate; (C) is the change in the heat of fusion with heating rate. Source: Reproduced from Wang, C., Tsou, S.-Y, Lin, H.-S., 2012. Colloid Polym. Sci. 290, 1799 (Fig. 10) with permission from Springer Publishing.

Thermal analysis of aliphatic nylon fibers

239

phases, and the melting point increases with heating rate for the HMT phase is most likely due to superheating. The HMT phase is thought to be highly oriented chains located in the skin layer that resulted from shear stresses during electrospinning.

13.4

Nylon-6,6

For nylon-6,6 the α phase, which has a trigonal crystal form, is generally formed from the melt and is the most thermodynamic crystal modification (Perkins and Porter, 1977). A Brill transition exists at high temperatures but below the melting point where a transformation from the α form to pseudohexagonal structure occurs (Wolanov et al., 2009). Reorganizational melting behavior in nylon-6,6 appears to have been observed first by White (1955) where he noticed a doublet melting peak in a DTA trace. The undrawn nylon-6,6 showed one peak from 240
C to 265
C, while the drawn fiber (with approximately 4 3 the extension) showed two distinct peaks from 240
C to 257
C and 259
C to 265
C. This was confirmed and further extended by Hybart and Platt (1967) who showed that this effect could be noted for different thermal treatments, such as annealed and precipitated samples, although as we discussed earlier in this chapter, the origin of the doublet peaks may be due to different causations of crystal perfection. Variable heating rate experiments were performed on fresh specimens of nylon-6,6 yarn by Sweet and Bell (1972), and they showed that the lower temperature peak grew at the expense of the upper temperature peak as the heating rate increased, typical of reorganization. DSC microscopy was used to determine when a nylon-6,6 fiber lost an imposed crimp. The fiber was extruded and drawn with a mechanical crimp introduced under heat and temperature. The crimp starts to relax prior to the complete crystalline melting with the left side of the y formed by the fiber, seen in Fig. 13.11, starting to open at the first melting peak, thus indicating that the first peak is related to the induced stress that formed the crimp. Bell and Murayama (1969) noted three damping peaks in DMA from 2180
C to 250
C in nylon-6,6. The α relaxation, which is related to the glass transition, occurs at about 100
C, while the β and γ relaxations are near 250
C and 2120
C, respectively. As with nylon-6, these relaxations are sensitive to both solvent and moisture. Dumbleton and Murayama (1970) studied the effect of drawing by DMA. They noted that the storage modulus (E0 ) curves for samples parallel and perpendicular to the stretch direction (drawn to a draw ratio of 3 3 ) do not cross with E0 parallel, which is always greater than E0 perpendicular, an indication of anisotropy. At a draw ratio of 4.5 3 , the nylon-6,6 curves were found to cross, which was attributed to the annealing necessary to achieve the higher stretch ratio. Dumbleton and Murayama (1970) also found that the α relaxation was only slightly affected by annealing but increasing orientation results in a higher temperature of the α transition.

240

Thermal Analysis of Textiles and Fibers

Figure 13.11 Several strands of nylon-6,6 carpet fiber as they are heated through their melting. The bend or crimp in the fiber straightens between 255
C and 259
C before the fiber becomes clear or isotropic at 263
C. Source: Reproduced from Weddle, B., Zemo, M., Sauerbrunn, S., 2007. DSC-microscopy of fibers, optical disk. In: Proceedings of the 35th NATAS Conference (Fig. 3) with permission from NATAS.

Murayama and Silverman (1973) studied the effects of molecular mass on nylon-6,6 film and fiber. DMA found the tan δ peak temperature shifted to higher temperatures (from 88
C to 120
C) with a decrease in peak height as the molecular mass increased. Felty and Murayama (1981) performed both dynamic tensile and compression mechanical analyses of crimped nylon carpet yarns. The Tg, as denoted by the tan δ maximum, was found to be about 105
C. Although both the tensile and compression testing gave similar peak temperatures, the intensity values differed between the two techniques. The room temperature tan δ intensity for compression testing was found to increase with the number of crimps per inch, which they believed was due to the highly crimped fiber having a high fiber-to-fiber friction. Bell (1968) found that the dye diffusion coefficient, D, can be related to DMA loss modulus, Ev, a plot of log (DT/Do) versus log ðE000 =ET00 Þ was found to be linear with a slope of 4.5 where the subscripts T and 0 indicate the measured temperature and reference temperature, respectively. This suggests that the dye diffusion is controlled by the mobility of the amorphous chain segments. Thermal shrinkage of nylon-6,6 fibers was observed to proceed above 130
C by a two-stage process (Ribnick, 1969). They reported an increase in shrinkage up to

Thermal analysis of aliphatic nylon fibers

241

180
C at which point the shrinkage is slowed (Ribnick, 1969). The maximum in shrinkage was opined to be related to the onset of creep in the nylon-6,6 yarns under the influence of a load. Anton (1973) reported a high rate of shrinkage for nylon-6,6 up to 70
C. From 70
C to 115
C, a slower shrinkage rate was observed and then, similar to Ribnick (1969), a high rate up to a peak at 175
C, where the yarn started to elongate, was to be then followed by plastic flow at 234
C (see Fig. 13.12). The lower rate of shrinkage between 70
C and 115
C was attributed to water loss. Manich et al. (2008) used the similar methodology as well as DSC to study yarn that had been textured. Texturing in this context is to simulate the properties of natural staple yarns and to this end they used two different lab processes: false-twist and air-jet. The false-twist technique is where the fiber is mechanically cold-twisted and then set by a dry heat at different temperatures. The second process was by a wet air-jet treatment where, in essence, the fiber is mechanically deformed in a hot, humid atmosphere. Texturing was found to slightly increase and broaden the glass transition. A crystallinity increase with texturing temperature from 35% to 37% was noted for the air-jet, whereas the false-twist texturing resulted in only a 5% increase. To characterize the fiber shrinkage, a temperature Ts was determined, which was the TMA curve derivative maximum below the melting point. The Ts of the airjet textured filaments was 140
C, whereas the Ts for the false-twist textured filaments were approximately 50
C below the texture temperature. Guerrimni et al. (2009) made electrospun nylon-6,6 fibers at ambient temperature from a formic acid solution. DSC analysis gave two melting endotherms: one between 248
C and 258
C, and the other between 258
C and 267
C. These were attributed to reorganization. The nanofibers had a lower crystallinity ranging from 15% to 30% compared to textile fiber, which is typically 40%.

Figure 13.12 Shrinkage force of nylon-6,6 fiber when heated. Source: Reproduced from Anton, A., 1973. Text. Res. J. 43, 524 (Fig. 6) with permission from Text. Res. J.

242

13.5

Thermal Analysis of Textiles and Fibers

Other aliphatic nylon fibers

As mentioned earlier, nylon-6 and nylon-6,6 account for most nylon commercial fiber production; some other nylons have been studied as fibers. Of academic interest is nylon-11 due to its piezo-activity; however, this activity is much less than polyvinylidene fluoride, a common piezoelectrical polymer (Mathur et al., 1984). Dosie`re (1989) investigated oriented fibers of nylon-11 annealed in contact with formic acid that after treatment could undergo large extensions. DSC analysis found that both the melting temperature and heat of fusion increased with raising treatment temperature, although it was noted that the lamellae became more regular when they reformed and they did so without increasing their thickness. Formic acid has also been used in the electrospinning of nylon-11 (Dhanalakshmi and Jog, 2008). This process was observed to produce the γ-form crystallites over the expected α-form that one would see from solution crystallization. The authors attributed this to a high elongation during electrospinning. When compared to the solution-crystallized control, the electrospun nanofiber from the solution showed an LMT and heat of fusion. The authors believe that a slower crystallization occurs for the solution casting versus a very fast crystallization due to the fast evaporation that occurs during electrospinning rather than the phase-change difference. Melt-spun nylon-12/carbon nanotubereinforced fibers have been studied by Sandler et al. (2004). DSC analyses were carried out at two different heating rates of 10 and 30
C/min. A glass transition temperature was measured at approximately 50
C, followed by some apparent cold crystallization, which appears constant and independent of nanofiller content. The onset of the main melting peak and peak temperature do not show any significant variations; however, for the faster heating rate, a lower temperature solder representing imperfect crystals is evident, but it disappears with the addition of the nanofiller. Fu and Chen (1983) studied the melting behavior of nylon-10,10 bristles where they noted multiple endothermic peak behavior that they attributed to structural reorganization. They were able to demonstrate an intervening exotherm for the typical dual reorganizational peaks by having a short isotherm just after the lower peak to fully allow its melting; after continuing heating, the exotherm was noted, followed by a final endotherm thus demonstrating that recrystallization is associated with the first melting. A nylon-6,6/6 (90/10) copolyamide cold-drawn yarn was studied by Anton (1973) besides nylon-6,6. He found that nylon-6,6/6 had a lower crystal perfection since nylon-6 and nylon-6,6 disrupt each other’s lattice packing, as well as a decrease in Tg. A terpolymer of nylon-6/6,6/10,10 (10/20/70) was electrospun from 2,2,2-trifluoroethanol (Li et al., 2006). As with the electrospinning of the nylon-11, reported earlier, the electrospun terpolymer showed a lower heat of fusion from that of the solvent cast due to more rapid crystallization. In conclusion, thermal analysis provides a valuable tool in the understanding of nylon fiber morphology. DSC provides information on the amorphous and crystalline phases from the typical analysis of glass transition and melting point to structural information on the reorganization of imperfect crystals. TGA can provide

Thermal analysis of aliphatic nylon fibers

243

decomposition and compositional information. TMA has been shown to offer elucidation toward understanding the complicated orientation that forms during the drawing process, while from DMA, besides obtaining mechanical property, measurements can detect subtle transitions.

References ABC-of-Fishing, 2015. ,http://www.abc-of-fishing.net/fishing-lines/types.asp. (23.05.15.). Acierno, D., La Mantia, F.P., Polizzotti, G., Alfonso, G.C., Ciferri, A., 1977. Polym. Lett. Ed. 15, 323. Anton, A., 1973. Text. Res. J. 43, 524. Arakawa, T., 1968. Polym. Lett. 6, 513. Arakawa, T., Nagatoshi, F., Arai, N., 1969a. Polym. Lett. 7, 115. Arakawa, T., Nagatoshi, F., Arai, N., 1969b. J. Polym. Sci.: A-2 7, 1461. Arimoto, H., Ishibashi, M., Hirai, M., 1965. J. Polym. Sci.: A 3, 317. Avramova, N., Fakirov, S., 1986. Polym. Sci.: B: Polym. Phys. 24, 761. Barriault, R.J., Gronholz, L.F., 1955. J. Polym. Sci. 18, 393. Bell, J.P., 1968. J. Appl. Polym. Sci. 12, 627. Bell, J.P., Murayama, T., 1969. J. Polym. Sci.: A-2 7 (7), 1059. Carrothers, W.H., 1938. U.S. Patent 2,130,523. Clark, Jr., R.L., 1996. Influence of the Interphase on the Mechanical Properties of Nylon 66 Composites (Ph.D. dissertation). Virginia Polytechnic Institute and State University. Dang, P., Tence-Girault, S., 2008. SKZ Polyamide Conference 3rd & 4th Dec 2008, Marienberg, Wurzburg, Germany. Dasgupta, S., Hammond, W.B., Goddard III, W.A., 1996. J. Am. Chem. Soc. 118, 12291. Davis, S., Grayson, R., Yalcinkaya, M., Wu, M., Yoneyama, M., 2014. Chemical Economics Handbook, Nylon Fibers, IHS Chemicals. Dennis, L.A., Buchanan, D.R., 1987. Text. Res. J. 57, 625. Dhanalakshmi, M., Jog, J.P., 2008. Polym. Lett. 2 (8), 540. Dole, M., Wunderlich, B., 1959. Makromol. Chem. 34, 29. Dolphin Sutures, 2015, ,http://www.dolphinsutures.com/nylon-sutures. (23.05.15.). Dosie`re, M., 1989. Macromol. Symp. 23, 205. Dumbleton, J.H., Murayama, T., 1970. Kolloid-Z. Z. Polym. 238 (12), 410. Electronic Code of Federal Regulations, 2014. Updated as of December 30, 2014 ,http:// www.ecfr.gov.. Felty, D.C., Murayama, T., 1981. J. Appl. Polym. Sci. 26, 987. Fu, S., Chen, T., 1983. Polym. Commun. 2, 145. Fu, Y., Annis, B., Boller, A., Jin, Y., Wunderlich, B., 1994. J. Polym. Sci.: B: Polym. Phys. 32, 2289. Gianchandani, J., Spruiell, J.E., Clark, E.S., 1982. J. Appl. Polym. Sci. 27, 3527. Greco, R., Nicolais, L., 1976. Polymer 17, 1049. Guerrini, L.M., Branciforti, M.C., Canova, M.C., Bretasa, R.E.S., 2009. Mater. Res. 12 (2), 181. Gupta, A., Saquing, C.D., Afshari, M., Tonelli, A.E., Khan, S.A., Kotek, R., 2009. Macromolecules 42, 709. Hybart, F.J., Platt, J.D., 1967. J. Appl. Polym. Sci. 11, 1449. Illers, K.H., 1977. Polymer 18, 551.

244

Thermal Analysis of Textiles and Fibers

Jaffe, M., Menczel, J.D., Bessey, W., 1997. Chapter 7: Fibers. In: Turi, E. (Ed.), Thermal Characterization of Polymeric Materials. Academic Press. Khanna, Y.P., 1990. J. Appl. Polym. Sci. 40, 569. Koyama, K., Suryadevara, J., Spuriell, J.E., 1986. J. Appl. Polym. Sci. 31, 2203. Kunugi, T., Suzuki, I., Akiyama, I., Hashimoto, M., 1982a. Polymer 23, 1193. Kunugi, T., Akiyama, I., Hashimoto, M., 1982b. Polymer 23, 1199. Kunugi, T., Ikuta, T., Hashimoto, M., Matsuzaki, K., 1982c. Polymer 23, 1983. La Mantia, F.P., Acierno, D., 1981. Colloid Polym. Sci. 259, 693. La Mantia, F.P., Polizzotti, G., Titomanlio, G., Accierno, D., 1982. J. Macomol. Sci. Phys. 21, 131. La Mantia, F.P., Saiu, M., Valenza, A., Paci, M., Magagnini, P.L., 1990. Eur. Polym. J. 26 (3), 323. Levchik, S.V., Weil, E.D., Lewin, M., 1999. Polym. Int. 48, 532. Li, Y., Goddard III, W.A., 2002. Macromolecules 35, 8440. Li, Y., Huang, Z., Lu¨, Y., 2006. Eur. Polym. J. 42, 1696. Liu, Y., Cui, L., Guan, F., Gao, Y., Hedin, N.E., Zhu, L., et al., 2007. Macromolecules 40 (17), 6283. Manich, A.M., Maı´llo, J., Cayuela, D., Carilla, J., Ussman, M., Gace´n, J., 2008. Effect of the air-jet and the false-twist texturing processes on the thermomechanical behaviour of polyamide 6.6 yarns. Therm Anal Calorim 93, 921926. Mathur, S.C., Scheinbeim, J.I., Newman, B.A., 1984. Piezoelectric properties and ferroelectric hysteresis effects in uniaxially stretched nylon-11 films. J. Appl. Phys. 56, 2419. Microlab Experiment, http://www.pslc.ws/macrog/lab/lab01.htm Copyright ©1998Microlab. Miller, R.L., 1989. In: 3rd edition Brandup, J., Immergut, E.H. (Eds.), Polymer Handbook, 1989. John Wiley & Sons, New York, Section VI, 46. Miura, H., Hirschinger, J., English, A.D., 1990. Macromolecules 23 (8), 2169. Miyata, S., Balakov, I.P., 1981. J. Macromol. Sci.-Chem. A16 (7), 1233. Murayama, T., Silverman, B., 1973. J. Polym. Sci.: Polym. Phys. Ed. 11, 1873. Murthy, N.S., 2006. J. Polym. Sci., B 44, 1763. Murthy, N.S., Bray, R.G., Correale, S.T., Moore, R.A.F., 1995. Polymer 36 (20), 3863. Ong, E.S., Kim, Y., Williams, H.L., 1986. J. Appl. Polym. Sci. 31, 367. Pagilagan, R.U., 1995. Chemistry. In: Kohan, M., I. (Ed.), Nylon Plastics Handbook. Hanser, pp. 3367. PerkinElmer, 2015. Application Note. ,http://www.perkinelmer.com/CMSResources/ Images/44-141236APP_010033A_01_Characterization_of_Single_Fibers% 20_or_Forensic_Applications_Using_High_Speed_DSC.pdf. (23.05.15.). Perkins, W.G., Porter, R.S., 1977. J. Mater. Sci. 12, 2355. Rahbar, R.S., Mojtahedi, M.R.M., 2011. J. Eng. Fibers Fabrics 6 (2), 7. Ribnick, A., 1969. Text. Res. J. 39 (5), 428. Rilsan, 2015. Rilsans Polyamide 11 in Oil & Gas Off-Shore Fluids Compatibility Guide. ,http://www.hclfasteners.com/literature. (23.05.15.). Rong, S.-D., Williams, H.L., 1985. J. Appl. Polym. Sci. 30, 2575. Sandler, J.K.W., Pegel, S., Cadek, M., Gojny, F., van Es, M., Lohmar, J., et al., 2004. Polymer 45, 2001. Schlack, P., 1941. Preparation of Polyamides, U.S. Patent 2,241,321. Serpe, G., Chaupart, N., 1996. J. Polym. Sci.: B: Polym. Phys. 34, 2351. Sriram, P., Blanchard, P., Smith, K., Yamaguchi, Y., 2013. Chemical Economics Handbook, Nylon Engineering Resins, IHS Chemicals.

Thermal analysis of aliphatic nylon fibers

245

Starkweather Jr., H.W., 1995. Transitions and relaxations. In: Kohan, M.I. (Ed.), Nylon Plastics Handbook. Hanser, pp. 139149. Starkweather Jr., H.W., Jones, P.G.A., 1981. J. Polym. Sci. Polym. Phys. Ed. 19, 467. Starkweather Jr., H.W., Zoller, P., Jones, G.A., 1984. J. Polym. Sci. Polym. Phys. Ed. 22, 1615. Supaphol, P., Lin, J.-S., 2001. Polymer 42, 9617. Sweet, G.E., Bell, J.P., 1972. J. Polym. Sci.: A-2 10, 1273. Telen, L., Puyvelde, P.V., Goderis, B., 2016. Macromolecules 49, 876. Todoki, M., Kawaguchi, T., 1977a. J. Polym. Sci. Polym. Phys. Ed. 15, 1507. Todoki, M., Kawaguchi, T., 1977b. J. Polym. Sci. Polym. Phys. Ed. 15, 1067. Tyre Cord, 2015. ,http://textilelibrary.weebly.com/uploads/1/1/7/4/11749432/tyre_cord_srk.ppt. (23.05.15.). Vinken, E., 2008. Polyamides: Hydrogen Bonding, the Brill Transition, and Superheated Water (Thesis). Eindhoven University of Technology. Wang, C., Tsou, S.-Y., Lin, H.-S., 2012. Colloid Polym. Sci. 290, 1799. Weddle, B., Zemo, M., Sauerbrunn, S., 2007. DSC-microscopy of fibers, optical disk. In: Proceedings of the 35th NATAS Conference. White, T.R., 1955. Nature (London) 175, 895. Wilhoit, R.C., Dole, M., 1953. J. Phys. Chem. 57, 14. Wolanov, Y., Feldman, A.Y., Harel, H., Marom, G., 2009. Polym. Lett. 3 (7), 452. Wunderlich, B., 1973. Macromol. Phys., vol. 1. Academic Press. Wunderlich, B., 1980. Macromol. Phys., vol. 3. Academic Press. Xenopoulus, A., 1990. Thermal Analysis and Studies of Conformational Disorder in Aliphatic Polyamides (Dissertation). Rensselaer Polytechnic Institute. Zachariades, A.E., Porter, R.S., 1979. J. Polym. Sci. Polym. Lett. Ed. 17, 277.

Thermal analysis of poly(aryl ether ketone) fibers

14

Lawrence Judovits Arkema, King of Prussia, PA, United States

Abstract Poly(aryl ether ketone), which is also denoted by polyaryletherketone (PAEK), is a family of linear aromatic polyether ketones. Two of the most well-known of the PAEK family are poly(ether ether-ketone) (PEEK) and poly(ether-ketone-ketone) (PEKK). PEKK is different from the other polyketones since it has a range of melting points due to the different isomeric ratios of its isophthalates and terephthalates (meta/para) linkages. Fibers made of both PEEK and PEKK exhibit a range of crystallinities with PEEK having a glass transition temperature of approximately 144 C with PEKK having a glass transition temperature of over 10 C higher. The melting point for commercial PEEK has been observed up to 340 C, while a melting point of up to 360 C has been reported for a commercial PEKK resin with a high terephthalate content. An equilibrium melting point of 395 C has been reported for PEEK, while a value of 410 C was published for PEKK with an allterephthalate form. Both PEEK and PEKK have high thermal stabilities with no significant decomposition until above 500 C. The main commercial resin producers of PEEK are Victrex, Solvay, and Evonik, while for PEKK it is Arkema and Solvay.

14.1

Introduction

Poly(aryl ether ketone), or more commonly called polyaryletherketone (PAEK), is a family of linear aromatic polyether ketones also referred to as polyketones. This resin family includes poly(ether ether ketone) (PEEK), poly(ether ketone) (PEK), and poly(ether ketone ketone) (PEKK), which are PAEK materials that have been reported to be made into fibers (see Table 14.1 for their structures). Other members of this family, mentioned in the literature, include poly(ether ether ketone ketone), poly(ether ether ketone ether ketone), and poly(ether ketone ether ketone ketone) though there is a lack of fiber data; hence, these resins are not covered in this chapter. In general, PAEK resins are thermoplastic resins with the ratio of the ether-toketone linkage dictating its properties, for example, the flexibility increases as the number of ether linkages increases except for mixed isomers, especially for the PEKK case, which will be covered later in this chapter. The most well-known of these is PEEK, both in general and for fiber applications, although there are some investigations on PEKK resin.

Thermal Analysis of Textiles and Fibers. DOI: https://doi.org/10.1016/B978-0-08-100572-9.00014-8 © 2020 Elsevier Ltd. All rights reserved.

248

Thermal Analysis of Textiles and Fibers

Table 14.1 Structures of poly(aryl ether ketone) resins used in fibers. Polymer

Structure

% of ketone linkages to bridge linkages

PEEK

33

PEKK

67

PEK

50

14.2

Poly(aryl ether ketone) synthesis

The two main synthetic routes in the polymerization of PAEK is either nucleophilic displacement or electrophilic polycondensation (Gan et al., 2001). Nucleophilic substitution process can be performed by causing the reaction of a dihydric phenol, such as hydroquinone, with a dihalo-benzoid compound under the presence of a base resulting in ether linkages (Attwood et al., 1981). Electrophilic polycondensation is carried out through a Friedel Crafts reaction in the presence of Lewis acid catalysts resulting in carbonyl bridges (Bonner, 1962). An example would be the polymerization of diphenyl ether and isophthalic acid using hydrogen fluoride as a solvent with a boron trifluoride/hydrogen fluoride catalyst complex (Thornton, 1969). The electrophilic process has the disadvantage of being both costly and environmentally harsh (Ward and Nield, 1988). The use of expensive raw materials and high disposal costs for PEEK synthesis therefore impacts the pricing of the resin (Frost & Sullivan, 2007). PEKK specifically uses diphenyl ether and a mixture of terephthaloyl (T) and isophthaloyl (I) chlorides as monomers (Bonner, 1962). The synthesis can be carried out through a polycondensation of an intermediate, which is prepared with an excess of diphenyl ether and then further reacted with phthaloyl chlorides (Gardner et al., 1992). A range of melting points is possible for PEKK by adjusting the ratio of T to I isomers during polymerization (Gay and Brunette, 1989).

14.3

Producers

Major producers of PEEK are Victrex and Solvay with Victrex having the highest revenues reported for both North America and Western Europe in 2012 (Frost & Sullivan, 2013). In the same study, Solvay was deemed a distant second although it has expanded its PAEK line with the acquisition of the polymer division of Gharda Chemicals and Cytec (Frost & Sullivan, 2013; Cytec, 2015).

Thermal analysis of poly(aryl ether ketone) fibers

249

Table 14.2 Examples of commercial poly(aryl ether ketone) (PAEK) resin with reported data. Polymer

Manufacturer

Trade name

Tg ( C)

Tm ( C)

References

PEEK

Victrex

144

334

Victrex (2014)

PEEK

Evonik

No data

B340

Evonik (2011)

PEKK

Arkema

160

Amorphous

PEKK

Arkema

162

331 334

PEKK

Arkema

165

357 360

PEKK PEK

Cytec/Solvay Gharda/Solvay

Victrex PEEK 450G VESTAKEEP 4000G Kepstan 6000 series Kepstan 7000 series Kepstan 8000 series Cytec G-PAEK 1200P

159 152 154

337 370 372

Kepstan 6000 (2013) Kepstan 7000 (2012) Kepstan 8000 (2013) Cytec (2017) Songhan (2017)

Evonik was ranked third (Frost & Sullivan, 2013) but has also significantly expanded its production at its site in Changchun, China (Plastech, 2015). Manufacturers of PEKK are Arkema and Cytec (now part of Solvay) with Arkema offering a wide range of melting points for its PEKK Kepstan line (see Table 14.2). Zyex (2015) is a producer of monofilament and multifilament PEEK and may introduce other PAEK fibers. Zyex (2015) sells its fibers to aerospace, dry filtration and chemical separation, and medical markets. Their fibers can be woven into a range of textile structures as well sold as nonwovens, especially for conveyor belts (Zyex, 2015). Zeus (2017) also makes PEEK drawn fiber that is used in tube braiding. PEEK fiber is now under investigation to be used for a wide range of biomedical and orthopedic applications due to its range of structural and mechanical properties (Ouellette and Gibert, 2015).

14.4

Thermal analysis basics of poly(ether ether ketone) and poly(ether ketone ketone)

Most of the literature on the PAEK family is, for the most part, concentrated on PEEK with some investigations on PEKK. The equilibrium melting point for PEEK is given by Cheng et al. (1986) as 395 C with a 100% crystallinity heat of fusion of 124 J/g with a Tg of 146 C and a ΔCp of 0.27 J/(g C). Slightly different unit-cell parameters have been reported in the literature for PEEK. Two studies found that the unit cell was orthorhombic but with slightly different parameters (Dawson and Blundell, 1980; Hay et al., 1984). Dawson and Blundell (1980) reported unit-cell ˚ , b 5 5.86 A ˚ , and c 5 10.00 A ˚ , while a slightly different parameters of a 5 7.75 A

250

Thermal Analysis of Textiles and Fibers

˚ , b 5 5.922 A ˚ , and cell unit was reported by Hay et al. (1984) of a 5 7.781 A ˚ . In a systematic study of para-substituted PAEKs, Blundell and c 5 10.056 A ´˚ ´˚ Newton (1991) gave the unit cell of PEEK as a 5 7.76 6 0.15 A , b 5 5.89 6 0.1 A , ´˚ and c 5 9.95 6 0.1 A, the error of which encompasses (or very closely encompasses) both unit-cell parameters. This seems to be related to the conclusion by Hay et al. (1989), which says that the differences were due to a changing disorder in molecular packing with crystallization conditions and thus resulting in different crystalline densities. PEEK is susceptible to both cold crystallization and reorganization at typical differential scanning calorimetry (DSC) heating rates (Blundell and Osborn, 1983; Lee and Porter, 1987). Reorganization is where multiple peaks, seen in a DSC scan, represent a continuous melting and recrystallization where the imperfect crystals can perfect themselves during heating. To understand the crystallization of PEEK better, isothermal crystallization of PEEK was investigated by Cheng et al. (1986). Their investigation found three different crystallinities: a high-temperature melting fraction, which is usually the major fraction, a low-temperature melting fraction, and a fraction that formed on slow cooling from an isothermal hold. The crystallinity on cooling was determined by the difference in the total heat of fusion between samples analyzed with and without cooling. They found that the high-temperature melting peak formed first and was then followed by the low-temperature melting fraction through investigation of the analysis of different isothermal hold times at a crystallization temperature of 310 C. In addition, they found that increased heating rates increased the ratio of the low-to-high melting crystals since the low-melting crystals could reorganize. Cebe and Hong (1986) have also ascribed the upper temperature peak to form during isothermal crystallization. However, Blundell (1987) has theorized that the initial upper temperature peak formed immediately upon heating from a fast reorganization process and not directly from the isothermal crystallization. This was supported on samples that were initially crystallized at 210 C, postannealed at 250 C, 275 C, and 300 C and then analyzed at two heating rates: 20 and 80 C/min. He noted that both peaks were affected by increased heating rates with not only the lower temperature peak increasing in temperature, but also the upper temperature peak decreasing in temperature as typical of forming from reorganization. As mentioned earlier, the T/I or para/meta ratio has a profound effect on the PEKK melting point. The addition of the isophthalate or metaisomer, up to 50%, decreased the equilibrium melting temperature and increased the supercooling from a common starting melt temperature as can be seen in the Hoffman Weeks plots in Fig. 14.1 (Gardner et al., 1992). Hoffman Weeks plots are linear extrapolations of the melting point (Tm) from its isothermal crystallization temperature (Tc) to where it intercepts a Tm 5 Tc equilibrium line with no change in the crystal metastability before melting. Two orthorhombic crystalline forms were found. ˚, Form I was similar to that found for PEEK and has a unit cell of a 5 7.69 A ˚ b 5 6.06 A, and c 5 10.16, while form 2 has the unit-cell parameters of ˚ , b 5 5.75 A ˚ , c 5 10.16 A ˚ . Form 2 appeared after either cold crystallia 5 3.93 A zation or precipitation from solution, while form 1 was observed after

Thermal analysis of poly(aryl ether ketone) fibers

251

Tm (ºC)

400

350

PEKK (50/50) PEKK (60/40) PEKK (70/30) PEKK (80/20) PEKK (90/10) PEKK (100/0)

300

250 200

250

300 Tc (ºC)

350

400

Figure 14.1 PEKK Hoffman Weeks plots. Source: Reproduced from Gardner, K.H., Hsiao, B.S., Matheson Jr., R.R., Wood, B.A., 1992. Polymer 33 (12), 2483 (Fig. 12) with permission of Elsevier.

crystallization from the melt. Multiple melting peaks ascribed to reorganization can be seen for form 1 with double-melting peaks for the high terephthalate content and triple melting peaks for the high isophthalate content (up to 50%). Ho et al. (1994) determined a similar unit cell for form 1 although they obtained a larger form 2 unit cell that included two chains (in the unit cell) instead of one ˚ , b 5 11.3 A ˚ , and c 5 10.1 A ˚ . With the exception of with parameters a 5 4.2 A PEKK form 2, PEEK and PEKK form 1 are similar to the PEK unit cell of ˚ , b 5 6.00 A ˚ , and c 5 10.01 A ˚ (Blundell and Newton, 1991). The equia 5 7.76 A librium melting point for PEKK for all the para-linked isomers was approximately 410 C with a Tg of 165 C (Gardner et al., 1992). PEEK has a high thermal stability. Vasconcelos et al. (2014), through thermogravimetric analysis (TGA), found a two-step process with no significant decomposition until above 500 C in air and above 525 C under a nitrogen atmosphere as seen in Figs. 14.2 and 14.3. In nitrogen a residue of approximately 45% was measured at about 900 C resulting mainly from the loss of phenols. Mass spectrometry of the PEEK and PEK volatiles found hydroxyland hydro-terminated oligomers or phenol and benzofuran derivatives (Hay and Kemmish, 1987). From the decomposition fragments, Hay and Kemmish (1987) determined homolytic scissions of either the bridge ether or carbonyl bonds to produce radicals.

14.5

Poly(ether ether ketone) fiber

PEEK was melt-spun into a fiber and then hot-drawn by Zhen et al. (1985). The spinning and drawing temperatures were preliminarily estimated from the thermal

252

Thermal Analysis of Textiles and Fibers

110 100 Weight (%)

90 80 70 60 50 40 200

300

400 500 600 700 Temperature (ºC)

800

5ºC/min

10ºC/min

15ºC/min

20ºC/min

900

Figure 14.2 TGA of PEEK was performed under a nitrogen atmosphere at various heating rates using a Seiko EXSTAR6000. Source: Reproduced from Vasconcelos, G.C., Mazur, R.L., Ribeiro, B., Botelho, E.C., Costa, M.L., 2014. Mater. Res. 17 (1), 227 (Fig. 1) with permission of Mat. Res. 110

Weight (%)

90 70 50 30 10 –10 200

300

400 500 600 700 800 Temperature (ºC) 10ºC/min 5ºC/min 20ºC/min 15ºC/min

900

Figure 14.3 TGA of PEEK as performed under an air atmosphere at various heating rates using a Seiko EXSTAR6000. Source: Reproduced from Vasconcelos, G.C., Mazur, R.L., Ribeiro, B., Botelho, E.C., Costa, M.L., 2014. Mater. Res. 17 (1), 227 (Fig. 7) with permission of Mat. Res.

analysis results of a film specimen that was compression-molded and quenched in cold water. From the film specimen, they determined that the Tg was 143 C (determined using the hysteresis peak temperature), the cold crystallization peak temperature was 173 C, and the melting peak temperature was 337 C. On cooling a single

Thermal analysis of poly(aryl ether ketone) fibers

253

crystallization peak was observed with a maximum at 280 C. TGA found thermal stability to be greater than 400 C even in air. Based on the thermal analysis results, the fiber was spun at 380 C, and a smooth and steady spinning was obtained. The maximum hot draw ratio for the as-spun fibers were found to increase with the draw temperature and ranged from 2.8 at 140 C to 3.5 at 270 C for the highest spin-stretched fiber ratio of 36. In addition, PEEK fibers were made over a range of take-up speeds, draw ratios from 1 to 4, and annealing heat treatments from 200 C to 325 C (Shimizu et al., 1987). The crystalline and amorphous orientations as well as the crystallinity and crystallite size were determined from wide angle X-ray diffraction (WAXD), small angle X-ray scattering (SAXS), density, and birefringence measurements. As-spun fibers with take-up velocities between 110 and 920 m/min were highly amorphous with crystallinities below 10% and showed cold crystallization by DSC at approximately 160 C when heated at 10 C/min using sample masses of 3 6 mg. This indicates that the as-spun fiber can be easily quenched to an amorphous state though cold crystallization was not observed after annealing. The amorphous region of asspun fibers was found to be more than ten times larger than found for the crystalline region, by birefringence measurements, suggesting that the crystallization occurred in the oriented region during spinning. The amorphous orientation showed a maximum at 250 C when annealed after drawing although if constrained at a fixed length, the maximum increased to 325 C. The Tg was found to increase linearly with the amorphous birefringence, indicating the effect of orientation on the amorphous phase. The initial modulus of the fibers before drawing was found to depend on the degree of orientation in the amorphous region. In general, Shimizu et al. (1987) suggest that the melt-spun fiber is close to amorphous and must be further processed to become crystalline. Song et al. (1989) studied the development of crystallinity and orientation in the melt-spun and then-drawn PEEK fibers. These fibers were drawn at 160 C, 180 C, and 230 C. Melt-spun fibers were characterized by DSC, birefringence, and WAXD. Their results built upon Shimizu et al. (1987), in which they found that at high drawdowns, the spinline stress can crystallize the PEEK fiber directly (Song et al., 1989). At low drawdown ratios (DDRs), the fibers are largely amorphous and exhibit significant cold crystallization. As the DDR increased, the crystallinity increased from 9% at a ratio of 57.7 to 21% for a ratio of 520. In general, DDR is calculated as the ratio of the linear take-up velocity to the extrusion velocity (Ouellette and Gibert, 2015). For the melt-spun fibers, the melting point was found to remain constant at approximately 335 C 340 C from a DDR of 10.4 1140 although the Tg was found to increase while the cold crystallization temperature decreased as analyzed by DSC for a heating rate of 20 C/min (see Fig. 14.4). This suggests that reorganization is not inhibited. The cold crystallization was found to completely disappear when the as-spun fiber was drawn. The increase in Tg with the DDR is due to the increased crystallinity inhibiting amorphous mobility. Young’s modulus and tensile strength increased with increasing spinline stress, while the elongation at break decreased. Ouellette and Gibert (2015) reported about melt-spun PEEK fibers and compared sample treatments of as-spun, after annealing, or hot-drawn fibers. The as-spun fibers

254

Thermal Analysis of Textiles and Fibers

Drawdown ratio

Heat flow exo →

10.4 32.6 57.7 520 760 1140

100

200

300

400

Temperature (ºC)

Figure 14.4 DSC of PEEK as-spun fibers heated at 20 C/min. Source: Reproduced from Song, S.S., White, J.L., Cakmak, M., 1989. Sen-i Gakkaishi 45 (6) 243 (Fig. 3) with permission of Sen-i Gakkaishi.

were produced using a custom-built spinner and were then posttreated with annealing at 280 C for 90 minutes or hot-drawn from 50% to 100% at 280 C. DDR of either 100 or 200 were used on two different molecular-mass grades, Evonik 4000G and Victrex 150G with the 150G being lower in molecular mass. DDR was specifically determined as the velocity at the take-up drum divided by the velocity of the molten resin calculated at the spinneret aperture. The fiber diameters ranged from 75 μm down to 30 μm for the hot-drawn fibers. The glass transition for the fingerprint reheats was barely detectable. For the initial heat, this was converse for the as-spun fibers, and their glass transition could be clearly seen. Cold crystallization was observable at 15 C 20 C above the Tg for the as-spun fiber, indicating low crystallinity for the melt-spun fiber. In addition, for both grades, a broad exotherm, with a peak temperature at approximately 280 C, was observed for both as-spun and hotdrawn fibers, indicating further recrystallization and melting. An annealing peak was observed at approximately 290 C for the annealed sample. For the 4000G a reheat Tg and melting peak of 147.5 C and 337.6 C, respectively, was observed, and the matching product literature is given in Table 14.2. A lower Tg of 143.5 C and a higher melting peak of 343.3 C were determined for the 150G sample. The lower molecular mass of the 150G would account for a lower temperature glass transition as well as a higher temperature melting peak due to greater reorganization.

14.6

Poly(ether ketone ketone) fiber

Judovits (2020) studied three PEKK monofilaments of different T/I ratios by DSC and TGA. All fibers had more than 50% terephthalate content. The results showed a melting point depression with increasing isophthalate content (or decreasing

Thermal analysis of poly(aryl ether ketone) fibers

255

terephthalate content) as well as a decrease in crystallinity for both the fiber and its reheat. All specimens were cycled at 10 C/min from ambient temperature to 400 C (see Figs. 14.5 and 14.6). For the initial heat, most of the crystallinity seen for the fiber is from cold crystallization that occurred during heating, which is probably due to quenching after the fiber was spun. Double-melting peak behavior was noted for the intermediate and high terephthalate content, which can be attributed to reorganization (Ho et al., 1994). Two types of DSC analysis were performed: unconstrained (or free-to-shrink conventional analysis) and constrained (or fixed length). The constrained analysis was performed by wrapping the fiber around a small aluminum disk having slits on opposite sides, tying the fiber ends, and then crimped in a standard or “classic” DSC pan. Only slight differences, primarily limited to the cold crystallization, were seen between the constrained and unconstrained specimens, which would indicate a low degree of orientation in the fibers. The reheats for each ratio matched for both the constrained and unconstrained crimping with the highest isophthalate content being amorphous and the highest terephthalate content having the highest melting point and heat of fusion. Finally, the glass transition temperature increased slightly with increasing terephthalate content. The TGA was performed with a 20 C/min heating rate on a TA Instruments 5000IR under an air or nitrogen atmosphere (see Figs. 14.7 and 14.8). TGA scans showed some noise near the glass transition, which was attributed to cold crystallization rather than shrinkage due to orientational effects. A high thermal stability was observed with a 5% mass loss above 545 C for the air runs and above

Figure 14.5 DSC of PEKK monofilament heated at 10 C/min with the fiber constrained and unconstrained. Source: From Judovits, L., 2020. To be published.

Figure 14.6 DSC reheats of PEKK monofilament cooled and heated at 10 C/min with the fiber constrained and unconstrained. Source: From Judovits, L., 2020. To be published.

Figure 14.7 TGA of PEKK monofilaments was performed under a nitrogen atmosphere at a 20 C/min heating rate using a TA Instruments Q5000IR. Source: From Judovits, L., 2020. To be published.

Thermal analysis of poly(aryl ether ketone) fibers

257

Figure 14.8 TGA of PEKK monofilaments was performed under an air atmosphere at a 20 C/min heating rate using a TA Instruments Q5000IR. Source: From Judovits, L., 2020. To be published.

555 C for the nitrogen runs. Also, for the nitrogen runs, a residue was measured at 900 C between 54% and 63%.

14.7

Conclusion

PAEK is a family of high-performance thermoplastics having both high glass transition and melting temperatures. For this family of resins, most literature, in general, and, on fibers, is limited to PEEK, which dominates the PAEK market although PEKK has been studied by some researchers. The properties of the PEKK resin are heavily influenced by the ratio of its para or terephthalate to meta or isophthalate isomer content.

References Attwood, T.E., Dawson, P.C., Freeman, J.L., Hoy, L.R.J., Rose, J.B., Staniland, P.A., 1981. Polymer 22, 1096. Blundell, D.J., 1987. Polymer 28, 2248. Blundell, D.J., Osborn, B.N., 1983. Polymer 24, 953. Blundell, D.J., Newton, A.B., 1991. Polymer 32 (2), 308. Bonner Jr., W.H., 1962. US Patent 3065205. Cebe, P., Hong, S.-D., 1986. Polymer 27, 1183.

258

Thermal Analysis of Textiles and Fibers

Cheng, S.Z.D., Cao, M.-Y., Wunderlich, B., 1986. Macromolecules 19, 1868. Cytec, 2015. ,https://www.solvay.com/en/press-release/solvay-successfully-completesacquisition-cytec-and-launches-integration-plans. (09.12.15.). Cytec, 2017. ,https://www.cytec.com/sites/default/files/datasheets/PEKK_032012.pdf. (04.01.17.). Dawson, P.C., Blundell, D.J., 1980. Polymer 21, 577. Evonik, 2011. ,http://industrial.vestakeep.com/sites/lists/RE/DocumentsHP/PI-VESTAKEEP4000G-EN.pdf. (October 2011). Frost & Sullivan, 2007. European Market for Ketones in European Markets for High Performance Polymers. Frost & Sullivan, 2013. Analysis of the North American and Western European High Heat Thermoplastics Market. Gan, D., Lu, S., Wang, Z., 2001. Polym. Int. 50, 812. Gardner, K.H., Hsiao, B.S., Matheson Jr., R.R., Wood, B.A., 1992. Polymer 33 (12), 2483. Gay, F., Brunette, C.M., 1989. US Patent 4816556. Hay, J.N., Kemmish, D.J., 1987. Polymer 28, 2047. Hay, J.N., Kemmish, D.J., Langford, J.I., Rae, A.I.M., 1984. Polym. Commun. 25 (6), 175. Hay, J.N., Langford, J.I., Lloyd, J.R., 1989. Polymer 30, 489. Ho, R.-M., Cheng, S.Z.D., Hsiao, B.S., Gardner, K.H., 1994. Macromolecules 27 (8), 2136. Judovits, L., 2020. To be published. Kepstan 7000, 2012. ,http://www.arkema.com/export/shared/.content/media/downloads/products-documentations/incubator/arkema-kepstan-7000-tds.pdf. (November 2012). Kepstan 6000, 2013. ,http://www.arkema.com/export/shared/.content/media/downloads/products-documentations/incubator/arkema-kepstan-6000-tds.pdf. (March 2013). Kepstan 8000, 2013. ,http://www.arkema.com/export/shared/.content/media/downloads/products-documentations/incubator/arkema-kepstan-8000-tds.pdf. (May 2013). Lee, Y., Porter, R.S., 1987. Macromolecules 20, 1336. Ouellette, E.S., Gibert, J.L., 2015. Polymer 63, 10. Plastech, 2015. ,https://www.plastech.biz/en/news/Evonik-expands-production-capacity-forPEEK-10064.. Shimizu, J., Kikutani, T., Ookoski, Y., 1987. Sen-i Gakkaishi 43 (10), 507. Song, S.S., White, J.L., Cakmak, M., 1989. Sen-i Gakkaishi 45 (6), 243. Songhan, 2017. ,http://www.lookpolymers.com/pdf/Gharda-Chemicals-G-PAEK-1200PUltra-High-Performance-PEK-Polymer.pdf. (04.01.17.). Thornton, R.L., 1969. US Patent 3,442,857. Vasconcelos, G.C., Mazur, R.L., Ribeiro, B., Botelho, E.C., Costa, M.L., 2014. Mater. Res. 17 (1), 227. Victrex, 2014. ,https://www.victrex.com/B/media/datasheets/victrex_tds_450g.pdf. (July 2014). Ward, M.V., Nield, E., 1988. US Patent 4,722,980. Zeus, 2017. ,https://www.zeusinc.com/materials/peek. (04.01.17.). Zhen, X.J., Kitao, T., Kimura, Y., Taniguchi, I., 1985. Sen-i Gakkaishi 41 (1), 59. Zyex, 2015. ,http://www.swicofil.com/zyex.html..

Surgical sutures

15

Joseph D. Menczel Thermal Measurements LLC, Fort Worth, TX, United States

Abstract The use and properties of absorbable (biodegradable: polyglycolic acid, polylactic acid, polydioxanone, silk, cotton, etc.) and nonabsorbable [nylons, poly(ethylene terephthalate) (PET), poly(tetrafluoro ethylene) (PTFE), poly(vinylidene fluride) (PVDF), polypropylene (PP)] sutures are explained in this chapter. Thermal analysis was used to differentiate between melting of crystal forms of poly(L-lactic acid) (PLLA). Differential scanning calorimetry (DSC) of high-strength solution-spun PLLA fibers is described. Detailed thermal analysis studies were used to find correlation between in vitro degradation time and property changes of polydioxanone. Thermal analysis was also used to study hydrolytic degradation of poly(lactide-co-glycolide).

By definition a surgical suture is a medical device that holds the edges of a surgical incision or a wound. A suture is applied on the wound or surgical incision with a needle. Suture thread is manufactured from various materials. Originally, sutures were made from natural materials, such as silk or cotton. Modern sutures, in most cases, are made of synthetic fibers. They must hold the tissue securely, at the same time they must be flexible enough to be knotted. Also, sutures should be hypoallergenic. The mechanical properties of suture fibers are very important. Lin et al. (1988) compared a number of different sutures regarding tensile strength, the handling characteristics, knot-breaking strength, and the knot rundown force of Maxon, Polysorb, and Vicryl. Considering the fiber construction, sutures can be monofilament or braided. Monofilament sutures are made from polypropylene (PP), catgut, nylon, polyvinylidene fluoride (PVDF), stainless steel, poliglecaprone, and polydioxanone, whereas braided sutures can be made from PGA (polyglycolide), polyglactin 910, silk, and polyester. The knotting properties of these two types of sutures are different. Sutures, mostly braided sutures, can also be coated. These include silk and polyester, braided or twisted nylon, poliglecaprone and polydioxanone sutures, PGA sutures, Catgut Chromic, and polyglactin 910. Sutures can be absorbable (biodegradable in the body) or nonabsorbable. Absorbable sutures are made of poly(glycolic acid) (PGA), poly(lactic acid) (PLA), polydioxanone, poly(trimethylene carbonate) and poly(ε-caprolactone), and also the previously mentioned copolymers and sometimes terpolymers. Absorbable sutures are used for internal body tissues. These sutures hold the body tissues together for sufficient time to allow Thermal Analysis of Textiles and Fibers. DOI: https://doi.org/10.1016/B978-0-08-100572-9.00015-X © 2020 Elsevier Ltd. All rights reserved.

260

Thermal Analysis of Textiles and Fibers

healing, but then the polymer will fully degrade. A major requirement is that the degraded products should not adversely interact with the body cells. For example, aliphatic polyesters, derived from aliphatic α-hydroxycarboxylic acid, biodegrade with the excretion of water and carbon dioxide resulting in no toxic consequences. Depending on the application, the diameter of synthetic absorbable sutures ranges from 0.02 to 0.7 mm. The necessary time for the biodegradation of sutures must be different, depending on the body part where they are used, from 3 5 days to 14 days. Hoffman (1977) gave a detailed summary on the medical use of polymeric fibers. Silk and cotton are also biodegradable, but their biodegradation is slow (Hoffman, 1977). Nonabsorbable sutures are made of nylon, poly(ethylene terephthalate) (PET), poly(tetrafluoro ethylene (PTFE), PVDF, PP, and stainless steel. They are used on skin wounds, from which they are removed after a certain time, or in stressful internal environments where absorbable sutures are not used due to their weaker mechanical properties. When cosmetic consequences are important, nonabsorbable sutures should be used because they cause less scarring due to smaller immune response. These sutures may be removed after a certain time, or left permanently on. Nonabsorbable sutures have trade names, such as Prolene (polypropylene), Ethilon (nylon) (Ethicon, Somerville, NJ). Sutures can also be classified on the basis of the fiber structure as monofilaments or multifilaments (or braided) according to their usage, such as orthopedic, general, cardiovascular, and ophthalmic sutures. The knot security of braided sutures is better, but better passage through tissues is provided by monofilament sutures. Also, lower tissue reaction is resulting from the use of monofilament sutures. Braided tissues sometimes have surface coating (such as wax, PTFE, and calcium stearate) to enhance properties, such as reduced tissue reaction, easier knotting ability, and easier passage through the tissues. The most attractive biodegradable sutures are the group of polyesters derived from aliphatic α-hydroxycarboxylic acid, because their biodegradation products have no toxic consequences: water and carbon dioxide are formed. The best known example of this is polydioxanone [ (CO CH2 O CH2 CH2 O)n] marketed by Ethicon, Somerville, NJ (PDS II): this is a monofilament suture with a melting point of 110 C. Another such suture, named VICRYL Plus, is antibacterial (polyglactin 910, a 90% glycolide (GA), and 10% L-lactide copolymer), whereas MONOCRYL, a copolymer from GA and ε-caprolactone, is bioabsorbable. Chromic Gut is absorbable but not synthetic (it is mostly collagen). All these sutures are the products of Ethicon. Sutures are monofilament synthetic absorbable surgical products prepared from a copolymer of GA and ε-caprolactone. Davis and Geck Co. from Danbury, CT markets the Dexon suture (PGA). Aesculap (B. Braun Surgical) markets MonoPlus (polydioxanone), Monosyn (a terpolymer from GA, trimethylene carbonate, and ε-caprolactone), Novosyn [poly(glycolide-co-L-lactide) 90/10], and Safil (polyglycolide). Kulkarni et al. (1971) studied poly(glycolic acid-co-L-lactic acid) copolymers and observed that the biodegradation of copolymers is faster than that of any of the PLA homopolymers (Frazza and Schmitt, 1971).

Surgical sutures

261

Gilding and Reed (1979) published a detailed paper about the copolymerization of glycolic and lactic acids by the ionic polymerization of the cyclic diesters of GA and L-lactide (LLA) (the catalyst was a tin-containing compound). In their prepared samples the Tg values of PGA and poly(LLA) (PLLA) were 36 C and 57 C, respectively (the real Tg values must be higher somewhat because the mentioned authors used the extrapolated onset as a transition temperature). Both PLLA and PGA are semicrystalline, and they can be easily quenched to the fully amorphous state. These authors also presented the phase diagram of the GA LLA copolymers: when the GA contents of these copolymers are between 25 and 70 mol.%, the copolymers are amorphous. Because of the mechanical integrity, the sutures must be semicrystalline; therefore they should have a composition of less than 25 or more than 70 mol.% GA. Thermogravimetric analysis (TGA) was very useful for the characterization of these polymers showing that the PGA homopolymer is the most thermally stable, followed by PLLA. No considerable degradation takes place in any of these homopolymers during melt processing, but some weight loss can be observed at the melting point: the degradation products are GA and/or lactide forming through a simple hydrolysis. Similar to several other polymers, unzipping degradation takes place when these polymers are irradiated with γ-rays for sterilization purposes. Hoogesten et al. (1990) found considerable differences in the melting behavior of PLLA fibers containing α and β crystal forms of PLLA (see Fig. 15.1). They found the melting curve of the fibers containing the mixture of α and β crystal forms to be the superposition of the melting curves containing α- and β-forms only. Naturally, the melting curve recorded in the constrained state was higher in temperature than that recorded in free-to-shrink state. The vast majority of publications on these fibers are coming from the group of Pennings. They carried out a pioneering work on thermal analysis of these fibers, similar to solution-spun polyethylene fibers. The authors studied preparation and properties of high-strength poly(L-lactide) fibers in detail (Eling et al., 1982; Gogolewski and Pennings, 1982, 1983a,b; Leenslag et al., 1984; Leenslag and Pennings, 1987a,b; Hoogesten et al., 1990). They prepared the fibers by both melt and solution spinning from several different solvents. As expected, the solution-spun fibers had better tensile properties. Gogolewski and Pennings (1983b) prepared PLLA fibers from a 7% solution in trichloromethane. Similarly, to the polyethylene fibers, the as-spun PLLA fiber had a broad melting endotherm due to the melting of folded chain lamellae, but the melting appeared as a very sharp peak for the hot-drawn fibers. Gogolweski and Penning observed a broad high-temperature tail on the melting peaks of the hot-drawn fibers (Fig. 15.2). This was attributed to “delayed randomization in the PLLA melt,” but it is not clear why the authors meant that. However, the fact that this hightemperature broad peak disappeared when the fiber was chopped to 2 3 mm pieces indicates that the high-temperature tail may correspond to the melting of extended chain crystals. It is possible that shrinkage was somehow retarded in the samples with longer fiber size. The heat of fusion increases considerably with an increasing draw ratio (DR).

262

Thermal Analysis of Textiles and Fibers

Endo

Heat flow (W/g)

(A)

c b a 160

180

200

190

210

Endo

Heat flow (W/g)

(B)

170

Temperature (ºC)

Figure 15.1 DSC melting curves of PLLA fibers: (A) curve a: melting of unconstrained PLLA fiber spun from a 4.5% solution in chloroform/toluene mixture, and then hot-drawn to DR 5 4 3 at 191 C. This fiber contains α crystal structure of PLLA only. Curve b: melting of unconstrained PLLA fiber spun from an 8% solution in chloroform, and then hot-drawn to DR 5 14 3 at 204 C. This fiber contains β crystal structure of PLLA only. Curve c: melting of unconstrained PLLA fiber spun from a 4% solution in 40/60 chloroform/toluene mixture, and then hot-drawn to DR 5 8 3 at 204 C. This fiber contains both α and β crystals of PLLA. (B) Constrained state measurement of the fiber shown in (A), curve a. PerkinElmer DSC-2 instrument, heating rate 5 C/min. Source: Reprinted from Hoogesten, W., Postema, A.R., Pennings, A.J., ten Brinke, G., Zugenmaier, P., 1990. Macromolecules 23, 634, Figs. 8 and 9 by permission of the American Chemical Society.

The drawing mechanism of PLLA was studied by Postema and Pennings (1989), and two appropriate temperature regions were found for melt drawing. One of these is below 180 C, and the deformation here takes place in the semicrystalline state, while the second one is between 180 C and 190 C, and the deformation proceeds in the melt phase. In this latter case, a semicrystalline state is obtained by strain hardening after the displacement of topological defects. An important conclusion was made about the necessity of applying low deformation rates to achieve homogeneous drawing. The tensile strength of such PLLA fibers can go up to 2.3 GPa.

Surgical sutures

263

Figure 15.2 DSC curves of poly(L-lactide) gel spun fibers. The fibers were spun from a 7% trichloromethane solution. Curve 1: as-spun fiber; curve 2: hot-drawn fiber (draw temperature 200 C), DR 5 10; curve 3: hot-drawn fiber (draw temperature 200 C), DR 5 16; curve 4: hot-drawn fiber (draw temperature 200 C), DR 5 20. PerkinElmer DSC-2 instrument, unconstrained measurements, 5 C/min. Source: From Gogolewski, S., Pennings, A.J., 1983b. J. Appl. Polym. Sci. 28, 1045, Fig. 10. Reprinted with permission of John Wiley and Sons.

Biodegradation of PLLA fibers is relatively slow under physiological conditions (Hoffman, 1977; Katz and Turner 1970; Reed and Gilding, 1981). This is the major reason for using the LLA GA copolymers for the preparation of real-life sutures. But Gogolewski and Pennings (1982) and Leenslag et al. (1984) found another solution to this problem: they prepared PLLA fibers with a highly loosened fibrillar structure by solution spinning from a good solvent (chloroform) in the presence of some additives (polyurethanes or camphor). Such a fiber structure has considerably high specific surface, and thus the body fluids can diffuse easily into the fiber ensuring faster degradation. The degradation rate was similar to the PGA PLLA

264

Thermal Analysis of Textiles and Fibers

fiber degradation. What was remarkable that the mechanical properties of such fibers were comparable to the earlier PLLA fibers. In all cases the crystallinity of the fibers increases considerably with increasing DR. The crystal-to-crystal transformation of PLLA was found by mechanical measurements (stress strain curves). The melting of these crystal forms is also different as can be seen in Fig. 15.1A. Various spinning conditions lead to various β/α crystal ratios in the fibers (the α form is helical, while the β-form in the chain has a planar zigzag conformation). Higher drawing temperatures tend to shift the structure of the fibers to have more β-form. Leenslag and Pennings (1987a) prepared high-strength PLLA fibers by dry spinning and subsequent hot drawing having a tenacity of 2.1 GPa and Young’s modulus of 16 GPa. The fibers were spun from chloroform/toluene mixtures at the Q conditions producing PLLA chains with an interrupted helical conformation. Of course, the ultimate tenacity of these fibers depends on the drawability and the cross section of the fiber. The drawability gets higher with increasing crystallinity of the as-spun fiber, and it also depends on the solvent composition, spinning conditions, and the concentration of PLLA in the solution. Yamshidi et al. (1988) measured the three most important thermal characteristics of PLLA, such as the glass-transition temperature, melting point, and the degradation behavior. Tg of the D- and L-isomers was essentially the same (57 C 58 C). The melting point for the crystalline L-isomer was 184 C. The authors used the number-average molecular mass, the melt viscosity, DSC and TGA data, and IRspectroscopic data to explain the degradation behavior of PLLA under various conditions. They determined that the stability of polylactides is sensitive to high temperatures ( . 190 C). According to the authors, the weakest point regarding the stability is the ester bond on the main chain. The thermal degradation products included water, monomers, oligomers, and polymerization catalysts. When the nonpolymeric contents were removed from the samples, and the hydroxyl end-groups were blocked, the thermal stability of these polymers improved. This group of scientists studied the thermal behavior of PLLA in detail and indicated that depolymerization of PLLA takes place beyond 200 C. They could suppress this depolymerization by acetylation of the chain ends of PLLA (replacing the OH-end groups). Penning et al. (1993) published a detailed paper describing the preparation of bioabsorbable suture fibers from copolymers of L-lactide with small amounts of Dlactide and ε-caprolactone (these two latter comonomers were in the fiber in the amount of less than 20%) in order to decrease the crystallinity: fibers with smaller crystallinity degrade faster. Every time there is an optimum content of the comonomer in the fiber: the fiber needs to be semicrystalline to ensure the necessary mechanical properties of the sutures, but they need to have as small crystallinity as possible to speed up the biodegradation. However, the mechanical properties are partially determined by the orientation of the amorphous phase. Penning et al. (1993) prepared fibers by dry spinning/hot drawing and compared these fibers to fibers prepared by melt spinning/hot drawing. The 85/15 LLA/DLA copolymer fiber was fully amorphous, and low orientation could be achieved in such a fiber. Its optimum draw temperature was at the glass-transition temperature. At the

Surgical sutures

265

optimum draw temperature in copolymers with other compositions, low crystallinity, could be achieved, and these fibers solidified slowly during dry spinning. Hot drawing did not raise the crystallinity of such fibers, and the orientation level was limited. Dry spinning produced a 1 GPa maximum tensile strength fibers, which is half the tensile strength of the PLLA fibers. The glass-transition temperature was found to be the most important parameter affecting the maximum achievable modulus of these copolymer fibers: the optimum draw temperature for low-crystallinity copolymers was 40 C 50 C higher than Tg. Melt spinning leads to a fiber with a maximum tensile strength of 0.3 GPa. Also, significant molecular mass decrease was observed in melt spinning. Thus this research group recommended that decreasing the degradation may lead to fibers with higher tensile strength and modulus. As could be expected, due to its importance, most publications dealt with the degradation of sutures. Most sutures are biodegradable by hydrolytic degradation. What concerns thermal analysis, several papers dealt with the correlation of hydrolytic degradation with crystallization ability of suture polymers. Ping and Cameron (2002) investigated the morphological changes in polydioxanone sutures (PDSII) with small angle X-ray scatterring (SAXS), wide angle X-ray scattering (WAXS), optical microscopy, DSC, and dynamic mechanical analysis (DMA). The initial hydration causes significant changes in the properties of the suture. The authors made measurements in the time range of 0 80 days and found two stages for the degradation. In the first stage, mainly chain scission takes place, and this stage starts immediately after the suture is hydrated, but the molecular mass decrease does not influence the physical properties yet. Then, there is a gradual transition to the second stage in which the changes in physical and mechanical properties become apparent. The borderline between the first and second changes is about 15 days. The morphological changes in polydioxanone sutures were studied by Ping and Cameron (2002). They utilized DSC, WAXS, SAXS, DMA, and mechanical measurements to find correlation between in vitro degradation time and the changes in properties. The sutures exhibited significant changes in properties on initial hydration. The time frame of studies was 0 80 days, and two subsequent stages of degradation were revealed. The chain scission takes place in the first stage (lasting about 15 days) and starts as soon as the chains become hydrated. Nevertheless, the effects of scission on the physical properties can hardly be seen at this stage; therefore this stage was called “dormant.” Then, the dormant stage gradually changes to the second (active) stage. In the active stage the changes in the physical structure and the mechanical properties of the suture can easily be measured. Sabino et al. (2004) also studied the hydrolytic degradation of high molecular mass poly(p-dioxanone) (PPDX). They degraded the samples in distilled water or in a phosphate buffer at body temperature (37 C). The degradation of PPDX takes place in two steps depending on the state of the macromolecules: the amorphous regions degrade faster as expected, whereas the degradation of crystalline regions is much slower. The changes in the sample structures were monitored by viscosity measurements, DSC, polarization optical microscopy, mass loss, pH measurements, and scanning electron microscopy. Spherulitic crystallization was observed by

266

Thermal Analysis of Textiles and Fibers

polarization microscopy. The spherulitic crystallization was similar for all samples, changes were observed below a certain critical molecular mass only. The lamellar thickness developed during isothermal crystallization was determined by atomic force microscopy and DSC melting-point measurements. The surface-free energy of PPDX was estimated by the Thomson Gibbs equation to be 166 erg/cm2 for the surface perpendicular to the chain direction, which is comparable to the values obtained previously for similar linear polyesters. Loo et al. (2005) described how isothermal annealing influences the hydrolytic degradation rate of poly(lactide-co-glycolide) (PLGA). They carried out isothermal crystallization followed by annealing at 115 C. The goal of these heat treatments was to increase the crystallinity of the polymer. 60 minute annealing was needed to achieve the maximum increase in crystallinity. As expected, the crystal size and/or perfection increased with increasing annealing time. These films then were hydrolytically degraded in phosphate-buffered saline solution (pH 7.4) at body temperature (37 C) for up to 150 days. The authors observed insignificant mass loss in the samples and concluded that this was still the first stage of the degradation. As mentioned previously, annealing increases the crystallinity, and this increased crystallinity retarded the hydrolytic degradation when compared to the unannealed samples. Schliecker et al. (2003) published a very interesting paper. Since oligomers were thought to accelerate the hydrolytic degradation of PLGA due to their increased number of carboxylic end groups, the authors decided to check this: two D,L-lactic acid oligomers with molecular masses close to their critical limit of solubility were synthesized and were incorporated into PLGA films. Three different concentrations were used (0%, 10%, and 30% m/m). The initial films were amorphous. When the oligomer concentration of the samples increased, the glass-transition temperature (Tg) decreased. The results indicated that the initial mass loss and water absorption increase in oligomer-containing films. The increase is a function of the average molecular mass and the oligomer concentration, but the presence of oligomers did not accelerate the degradation. On the contrary, the lag time to the start of the polymer erosion increases in the presence of oligomers and this was explained by higher crystallinity of the oligomer-containing samples. It was also found that the crystallization process of the oligomer-containing samples starts earlier, and it is clear that this is the consequence of a higher mobility of the oligomers. Blaker et al. (2005) used several thermal analysis techniques to characterize commercial sutures. Some of the sutures were prepared with antimicrobial coatings from silver-doped bioactive glass (AgBG) interlocking particulates, and their thermal properties were compared to sutures without the mentioned antimicrobial coatings. The glass they used had a general formula of SiO2 CaO P2O5 Na2O, and these glasses formed firm bonds with both hard and soft tissues in the human body. The authors investigated the effect of a slurry dipping technique: this technique is used to coat resorbable Vicryl (polyglactin 910) and nonresorbable Mersilk surgical sutures. The used analytical techniques were conventional and modulated temperature DSC. Differential thermal analysis (DTA) and thermogravimetric analysis (TGA) were applied to establish the temperatures of thermal degradation and to determine the quantity of the AgBG coatings on the suture surfaces. The slurry

Figure 15.3 DSC heating curves of undegraded and degraded polydioxanone filament. The left column shows the DSC curves of undegraded filaments, whereas the right column shows the DSC curves of degraded filaments. Time of degradation: 6 weeks. (A and B): annealing temperature: 65 C. The DSC curves from top to bottom: unannealed sample; annealed for 3 hours; annealed for 6 hours; annealed for 9 hours; annealed for 12 hours; annealed for 24 hours (C and D). Annealing temperature: 75 C; the DSC curves from top to bottom: unannealed sample; annealed for 3 hours; annealed for 6 hours; annealed for 9 hours; annealed for 12 hours; annealed for 24 hours (E and F): Annealing temperature: 85 C; the DSC curves from top to bottom: unannealed sample; annealed for 3 hours; annealed for 6 hours; annealed for 9 hours; annealed for 12 hours; annealed for 24 hours. TA Instruments Q2000 DSC, heating rate 10 C/min. Source: From Abhari, R.E., Mouthuy, P.A, Zargar, N., Brown, C., Carr, A., 2017. J. Mech. Behav. Biomed. Mater. 67, 127, Fig. 5.

268

Thermal Analysis of Textiles and Fibers

Figure 15.4 Heat of fusion of undegraded and degraded polydioxanone filament before (A) and after (B) annealing at various temperatures. Part (C) shows the heat of the annealing peak. Time of degradation: 6 weeks. Source: From Abhari, R.E., Mouthuy, P.A, Zargar, N., Brown, C., Carr, A., 2017. J. Mech. Behav. Biomed. Mater. 67, 127, Fig. 6.

dipping technique itself did not influence the thermal transitions in both types of the sutures. Weak transitions were measured using heating rates as high as 100 C/ min. The simultaneous DTA TGA measurements indicated a greater amount of AgBG on the Mersilk suture surfaces. Abhari et al. (2017) showed how thermal analysis can be applied to obtain practically important results. They annealed polydioxanone filaments at various temperatures in order to improve the mechanical properties of electrospun fibers. Undegraded and degraded filaments were studied. The DSC curves are shown in Fig. 15.3, and the heat of fusion values measured after various annealing times are shown in Fig. 15.4. Annealing leads to the development of a small melting peak at around 80 C as usual for semicrystalline polymers, but it caused marginal improvement only in the tensile strength. The crystallinity of these fibers did not really increase as can be seen in Fig. 15.4A. For the degraded samples the crystallinity of the not-annealed fibers was considerably higher indicating that the molecular mass breakdown (due to degradation) increases the crystallization ability of these materials, as expected. Annealing at 75 C and 85 C results in increased crystallinity for

Surgical sutures

269

the undegraded fibers, while it leads to decreased crystallinity in the case of degraded fibers. Perhaps, these phenomena could be explained if the thermal history of the samples after the annealing was known. Unfortunately, the authors did not report how the samples were cooled to room temperature after the annealing process. This can be critical when evaluating the melting behavior (ΔHf) of the annealed fibers. If the cooling in these experiments was uncontrolled, it would be difficult to compare the the results of the DSC curves and the mechanical measurements, because the crystallinity of the fibers can be different. The authors found that after 6 weeks of annealing at 65 C creates a structure that retains more tensile strength than the unannealed fibers. On the other hand, annealing at 75 C and 85 C has the opposite effect: the annealed samples lose their mechanical properties more than the unannealed samples. Also, it was found that the small endotherm due to the annealing slowly disappears after about 3 weeks. This indicates that the small crystallites developed by the annealing are unstable. It seems that both the melting behavior and the behavior of the amorphous phase (glass transition) should be recorded in this case, in order to make definite conclusions about the thermal behavior of the filaments.

References Abhari, R.E., Mouthuy, P.A., Zargar, N., Brown, C., Carr, A., 2017. J. Mech. Behav. Biomed. Mater. 67, 127. Blaker, J.J., Boccaccini, A.R., Nazhat, S.N., 2005. Biomater. Appl. 20 (1), 81. Eling, B., Gogolewski, S., Pennings, A.J., 1982. Polymer 23, 1987. Frazza, E.J., Schmitt, E.E., 1971. Biomed. Mater. Symp. 1, 43. Gilding, D.K., Reed, A.M., 1979. Polymer 20, 1459. Gogolewski, S., Pennings, A.J., 1982. Makromol. Chem. Rapid Commun. 3, 839. Gogolewski, S., Pennings, A.J., 1983a. Colloid Polym. Sci. 261, 477. Gogolewski, S., Pennings, A.J., 1983b. J. Appl. Polym. Sci. 28, 1045. Hoffman, A.S., 1977. J. Appl. Polym. Sci. Polym. Symp. 31, 313. Hoogesten, W., Postema, A.R., Pennings, A.J., ten Brinke, G., Zugenmaier, P., 1990. Macromolecules 23, 634. Katz, A.R., Turner, R.J., 1970. Surg. Gynecol. 131, 701. Kulkarni, R.K., Moore, E.G., Hegyeli, A.F., Leonard, F., 1971. J. Biomed. Mater. Res. 5 (3), 169. Leenslag, J.W., Pennings, A.J., 1987a. Polymer 28, 1695. Leenslag, J.W., Pennings, A.J., 1987b. Makromol. Chem. 188, 1809. Leenslag, J.W., Gogolewski, S., Pennings, A.J., 1984. J. Appl. Polym. Sci. 29, 2829. Lin, G.T., An, K.N., Amadio, P.C., Cooney, W.P., 1988. J. Hand Surg. Am. 13A, 553. Loo, S.C., Ping, O.C., Wee, S.H., Boey, Y.C., 2005. Biomaterials 26 (16), 2827. Penning, J.P., Dijkstra, H., Pennings, A.J., 1993. Polymer 34, 942. Ping, O.-C., Cameron, R.E., 2002. J. Biomed. Mater. Res. 63 (3), 280. Postema, A.R., Pennings, A.J., 1989. J. Appl. Polym. Sci. 37 (8), 2351. Reed, A.M., Gilding, D.K., 1981. Polymer 22, 494.

270

Thermal Analysis of Textiles and Fibers

Sabino, M.A., Albuerne, J., Mu¨ller, A.J., Brisson, J., Prud’homme, R.E., 2004. Biomacromolecules 5 (2), 358. Schliecker, G., Schmidt, C., Fuchs, S., Wombacher, R., Kissel, T., 2003. Int. J. Pharm. 266 (1-2), 39. Yamshidi, K., Hyon, S.-H., Ikada, Y., 1988. Polymer 29 (12), 2229.

Further reading Chu, C.-C., 2002. Chapter 5: biodegradable polymeric materials. In: Park, J.B., Bronzino, J. D. (Eds.), Biomaterials, Principles and Applications. CRC Press, ISBN: 978-0-84931491-9. Domb, A.J., Kost, J., Wiseman, D.M., 1997. Handbook of Biodegradable Polymers. Overseas Publishers Association, Amsterdam B.V.

Thermal analysis of poly(vinyl alcohol) fibers

16

Zohar Ophir New Jersey Innovation Institute, Newark, NJ, United States

Abstract These days the main uses of poly(vinyl alcohol) fibers are in industrial applications, where high strength, useful properties over wide range of temperatures, durability, solvent resistance, and adhesion to matrix materials play a major role. Thermal analysis methods, such as differential scanning calorimetry (DSC), dynamic mechanical analysis (DMA), and thermogravimetric analysis (TGA), of the fibers serve as useful tools, both for manufacturers and for end users, for optimization of the production process conditions, quality control, and assessing the ultimate performance in each application. DSC has been widely used for monitoring the development of morphology of the fibers, as they go through the different production stages. DMA and thermomechanical analysis tests provide information how the fibers morphology translates to useful mechanical properties over the temperature range of interest, and TGA serves mostly for determination of moisture content and thermal degradation of the fibers.

16.1

Introduction

The technology of poly(vinyl alcohol) (PVA) fibers was developed during the first half of the 20th century, mostly for textile applications (Sakurada, 1985), but these fibers were replaced later by other synthetic fibers, such as nylons and polyester. These days the main commercial uses of PVA fibers are in industrial fibers, of which the most important application is fiber reinforcement of cements (Magalheas et al., 2013; Uddin et al., 2011). The benefits of the addition of highstrength fibers into cements are toughening by prevention of crack propagation in the brittle cement matrix. The main properties that make PVA particularly attractive for reinforcing cement are the combination of low cost, high strength and high modulus, good endurance in alkaline environment, and excellent adhesion to cement because of large number of hydroxyl groups along the chain. These hydroxyl groups provide hydrogen-bonding sites (Sakurada, 1985; Yumin et al., 2009; Chen et al., 2007), which are major contributor to the adhesion strength of these fibers. Although PVA is water soluble, its fibers solubility in water can be reduced, or prevented, by heat treatment, which induces and perfects crystallization, and by cross-linking (Sakurada, 1985).

Thermal Analysis of Textiles and Fibers. DOI: https://doi.org/10.1016/B978-0-08-100572-9.00016-1 © 2020 Elsevier Ltd. All rights reserved.

272

Thermal Analysis of Textiles and Fibers

Other notable properties of PVA fibers are high crystallinity, abrasion resistance, environmental stability, chemical resistance to most solvents, oils and weather conditions, and melting temperature in the range of 220
C240
C (Sakurada, 1985).

16.2

Manufacturing of poly(vinyl alcohol) fibers

PVA fibers are usually manufactured from medium molecular weight grades of PVA (Sakurada, 1985) in which the degree of polymerization (DP) is about 1700 (average molecular mass of about 75,000). An intermediate material in the production process of PVA is polyvinyl acetate. The acetate groups are removed in the next stage by hydrolysis, leaving behind hydroxyl groups and different percentages of hydrolysis are used for obtaining desired properties of the final products. PVA for fiber manufacturing needs to have high degree of hydrolysis (98%99%) in order to improve its stability in hot water. Despite being a linear, flexible-chain polymer, PVA cannot be melt spun into fibers because it starts to degrade at temperatures close to its melting point (Sakurada, 1985), unless the melting point is suppressed by large amount plasticizers (Chen et al., 2007; Jang and Lee, 2003; Li et al., 2010a,b), but adding large amount of plasticizers to PVA significantly worsens some of its attractive properties (Wu et al., 2012). Most of the industrial processes for manufacturing PVA fibers are based on solutions of PVA in water, or in organic solvents (Sakurada, 1985), including wet, dry, and gel spinnings. Electrospinning of PVA nanofibers from solutions has been also demonstrated, mostly for biomedical applications (Shaikh et al., 2012).

16.2.1 Wet spinning of poly(vinyl alcohol) fibers Wet spinning of PVA from water solution is the primary manufacturing process of these fibers for typical industrial uses (Sakurada, 1985). The typical PVA grade for wet spinning is almost fully hydrolyzed and has a DP of about 1700. The concentration of the spinning solution is usually in the range of 14%16%, and the fibers are extruded through multihole spinnerets, which can have up to 60,000 holes. The coagulation is performed by passing the strands through high-concentration solution of salts, such as sodium sulfate, or concentrated NaOH, or KOH, hydroxide solution. In both cases the coagulation is performed by the extraction of water from the polymer dope to the concentrated bath solution. The coagulation is usually performed in two or more baths, after which the fibers are washed while being stretched in water bath. Stretching is carried out at these stages in order to increase the orientation of macromolecular segments while preventing the fiber to get dissolved again in the wash bath. Following washing, the fibers are passed through drying, hot drawing, heat setting, and stabilization stages. The hot drawing steps increase the orientation of the fibers and their crystallinity, while the primary objective of the heat setting is to reduce the solubility of the PVA fibers in cold or in warm water.

Thermal analysis of poly(vinyl alcohol) fibers

273

There are very few solvents for PVA, which can be used as alternative manufacturing methods (Sakurada, 1985). These solvents include dimethyl sulfoxide and several glycols. These spun fibers can be coagulated with acetone, methanol, toluene, or their mixtures. However, the technical difficulties arising during application of highly volatile or flammable liquids annul the resulting slight property improvement of these fibers (Sakurada, 1985).

16.2.2 Dry spinning of poly(vinyl alcohol) fibers Dry spinning of PVA fibers is performed with concentrated solution of PVA in water (Sakurada, 1985). In this process, water is removed by evaporation, making it a cleaner and more environmentally friendly process than wet spinning. The PVA concentration can be as high as 43% if the dope is kept at high temperatures (up to 160
C). This is different from the wet spinning process, which is performed at much lower temperatures and dope concentration. The properties of the PVA fibers that are produced by this method are good. However, the drawback of the dry spinning process is that the low number of the spinneret holes (about 200, compared many thousands of spinneret holes in wet spinning processes).

16.2.3 Melt spinning of poly(vinyl alcohol) fibers PVA itself cannot be melt spun because of its thermal degradation below the melting point. Nevertheless, melt spinning can be performed if the melting point is suppressed by addition of certain low-molecular mass materials. This can be achieved by adding several plasticizers, such as poly(ethylene glycol) and other polyols, poly (ethylene oxide), in some solvents, such as water. When the water concentration is in the order of 50%, the PVA/water solution is very viscose even at high temperatures, and it can be processed as a melt (Wu et al., 2010). Plasticization by water has some advantages. This is an environment-friendly process, and the water can easily be removed by evaporation. Another advantage of plasticizing by water is that it can be easily removed by evaporation at the end of the process, while conventional plasticizers remain blended with the PVA in the final product and reduce its mechanical and solvent resistant properties.

16.2.4 Gel spinning of poly(vinyl alcohol) fibers The development of gel spinning technology for manufacturing high-strength fibers from ultrahigh molecular weight (UHMW) polyethylene (PE) encouraged attempts to use this technology with other polymers. Having a similar zigzag molecular structure as PE, it was anticipated that PVA has also the potential for achieving high mechanical properties once its high molecular weight (MW) chains are fully oriented (Sakurada et al., 1967; Tashiro et al., 1977). Kwon et al. (1984) patented a process for gel spinning of UHMW PVA (2 3 106) fibers that were made from very pure vinyl acetate. The nascent fibers were drawn by a ratio of about 2 in the coagulation bath and then drawn at high temperatures (up to 154
C) by a factor of

274

Thermal Analysis of Textiles and Fibers

5. These fibers have tenacity of about 2 GPa and modulus of about 60 GPa, which are very respectable, but still significantly less than the properties of UHMW PE with DP of over 100,000. Attempts to exploit this potential were only partially successful, and the expectations of making PVA fibers with similar mechanical properties to those of gel spun, while having better high-temperature performance and good adhesion to matrix materials, were not fulfilled so far. Honywell’s Spectra 1000100 fibers have tenacity and modulus of 43 and 133 GPa, respectively. Cebe and Grubb (1985) managed to increase the hot drawing ratio of medium MW PVA fibers to 14 and measured modulus of 20 GPa. The inability to draw the fibers by a much larger factor, as done in gel spun UHMW PE, was attributed to the presence of hydrogen bonds, which create strong inter- and intrachain interactions. Garrett and Grubb (1986) managed to increase the draw ratio (Dr) to 38 and reported modulus of 60 GPa. Fibers, made from 180,000 molecular mass PVA (Mitchenko et al., 2011), had tensile strength and modulus of 2.2 GPa and 32.6 GPa, respectively. Kunugi et al. (1990) obtained strength and modulus of 2.4 and 70 GPa, respectively, after applying multiple drawing steps up to a Dr 15 and temperature of 260
C on gel spun 520,000 molecular mass PVA (DP of 11,800). Modulus of 64 GPa was obtained (Cha et al., 1994) from 220,000 molecular mass PVA (DP of 5000).

16.2.5 Hot drawing and stabilization of poly(vinyl alcohol) fibers Relatively low DRs can be achieved on the fiber prepared by wet spinning (Sakurada, 1985). This results in low-molecular orientation levels and results in poor fiber properties. Additional drawing is therefore needed for improving mechanical properties. The drawing can be performed in one stage with dry fibers or as a combination of wet and dry hot drawing stages. The full Dr, obtainable by these methods, can go up as high as 11, and the result is the increase in tenacity to about 8.8 cN/dtex (10 g/dtex) while reducing the elongation to 5%6%. Realizing that the existence of inter- and intramolecular hydrogen bonds may be limit the highest achievable Dr of the PVA fibers, several studies used plasticizers, such as poly(ethylene oxide), glycols, and amide containing groups (Chen et al., 2007; Li et al., 2010a,b), to increase the highest possible Dr and to improve the mechanical properties. Optimized drawing methods were developed for gel spun fibers, where the emphasis was on maximizing the Dr in order to increase the mechanical properties as much as possible. It was found that the drawing temperature needs to be very close to the melting point of PVA in order to soften the crystalline structure and overcome resistance of the hydrogen bond to molecular motions. Garrett and Grubb (1986) used a zone drawing method, where only a short section of the fibers is exposed at a time to the high temperature and then cooled rapidly, thus locking in the molecular orientation, while reducing reentanglement and thermal degradation. Kunugi et al. (1990) refined the zone drawing method by performing it in several stages and adding vibrations to the static tension. This way the crystalline structure was perfected gradually, and a very high drawing temperature of 260
C could be used in order to obtain Dr of 15.

Thermal analysis of poly(vinyl alcohol) fibers

275

Following the hot drawing process, heat treatment is used to increase the crystallinity of the fiber and to improve its hot water resistance (Sakurada, 1985). In certain cases, aldehydes are used in an additional step to introduce some crosslinking of the PVA macromolecules. The reason for cross-linking is to further improve of the stability of the PVA fiber in boiling water.

16.3

Thermal analysis of poly(vinyl alcohol) fibers

Several thermal analysis techniques have been used for the characterization of PVA fibers throughout all production stages from raw material to finished fibers. As a thermoplastic, semicrystalline polymer, the PVA feedstock for fibers manufacturing is usually tested by TGA (thermogravimetric analysis) and by DSC (differential scanning calorimetry). The main properties of interest at the stage of raw material are moisture content, decomposition temperature, melting temperature, and crystallinity. Being somewhat hygroscopic, PVA can absorb a few percent of water, which can be detected by TGA as weight loss from room temperature (RT) up to about 200
C (Gilman et al., 1994). The absorbed water can be detected both as a mass loss by TGA and simultaneously by DSC as an endothermic peak (indicating evaporation of water) in the same temperature range (Khasbiullin et al., 2014). Thermal degradation of highly hydrolyzed PVA (98%99%) starts above 220
C (Taghizadeh et al., 2015; Shie et al., 2002), while the melting temperature of various grades of PVA has been reported to be in the range of 220
C267
C (Sakurada, 1985), which explains why PVA cannot be melt spun into fiber, unless it is highly plasticized. The TGA curve in Fig. 16.1 shows that PVA loses 4.5% water between RT and 150
C. It can also be seen that thermal degradation starts at about 220
C. The glass transition temperature of PVA has been measured by several methods (Sakurada, 1985), and the reported values were in the range of 70
C85
C. The presence of crystallinity along with residual moisture makes the measurement of the glass transition difficult by DSC. The DSC curves in Fig. 16.2 show that the glass transition in the first heating of undried PVA powder was low (about 45
C), because the polymer was plasticized by water. The endothermic peak just above the glass transition temperature is attributed to moisture evaporation in the same temperature range as a mass loss step is observed by TGA in Fig. 16.1. The DSC curve of the same PVA (Fig. 16.2) following a 16 hours of drying process at 80
C in vacuum shows that the glass transition of the dry polymer is 91
C, its peak temperature of melting at 223
C, and its heat of fusion is 79 J/g. The approximate crystallinity of this sample was calculated from the DSC curve to be 57% by dividing the melting enthalpy of 79 J/g by the melting enthalpy 138.6 J/g of 100% PVA crystals (Peppas and Merrill, 1976). The effect of residual moisture in the fibers on their mechanical properties can be seen in Fig. 16.3 (Garrett and Grubb, 1986). The removal of residual water from

Filename: Operator ID: Sample ID: Sample Weight: Comment:

E:\Thermal Analy...\PVA 56-98 G2 powder.tgd Zohar PVA 56-98 G2 powder 6.018 mg

PerkinElmer thermal analysis

100

98

Water loss Delta Y = 4.379%

96

Thermal degradation 94

Weight %

92

90

88

86

84

82

80 20

50

150

100

200

250

300

320

Temperature (°C) 7/16/2015 12:58:52 PM 1) Heat from 25.00°C to 525.00°C at 10.00°C/min

Figure 16.1 TGA of undried MOWIOL 5698 PVA powder, obtained from Kuraray America. Source: Ophir, Z., Jaffe, M., 2014. Report to United Soybean Board, Contract Project 14405125286. Filename: Operator ID: Sample ID: Sample Weight: Comment:

C:\Program Fil...\PVA 56-98 G2 powder 2.dtd Zohar PVA 56-98 G2 powder 7.500 mg

PerkinElmer thermal analysis

–0.5 –0.4

Dried PVA powder

–0.2 0.0 Normalized heat flow endo down (W/g)

Tg: Half Cp extrapolated = 76.78°C

Delta Cp = 0.294 J/g °C

0.2 0.4 0.6

Area = 53.006 J/g Delta H = 53.0062 J/g

Undried PVA powder Tg: Half Cp extrapolated = 44.54 °C

0.8

Peak = 222

Delta Cp = 0.592 J/g °C

1.0

Evaporation of moisture

1.2

Area = 69.766 J/g Delta H = 69.7655 J/g

1.4 1.6 1.8

Peak = 223

2.0 0

20

40

60

80

100

120

140

160

180

200

220

240

Temperature (°C) 7/16/2015 12:37:14 PM 1) 2) 3) 4)

Hold for 2.0 min at -20.00°C Heat from -20.00°C to 240.00°C at 10.00°C/min Hold for 1.0 min at 240.00°C Cool from 240.00°C to -20.00°C at 10.00°C/min

5) 6) 7)

Hold for 2.0 min at -20.00°C Heat from -20.00°C to 240.00°C at 10.00°C/min Hold for 1.0 min at 240.00°C

Figure 16.2 DSC scans of not dried and dried PVA MOWIOL 5698 G2 powder (sample obtained from Kuraray America). Source: Ophir, Z., Jaffe, M., 2014. Report to United Soybean Board, Contract Project 14405125286.

Thermal analysis of poly(vinyl alcohol) fibers

277

50

0.4

40 λ = 22× = Dry = 13% H2O

0.3

20 0.2 10

tan δ

Storage modulus (GPa)

30

0 0.1 –10

–20 –100 –75 –50 –25 0 25 50 Temperature (ºC)

75

0.0 100

Figure 16.3 Storage modulus of PVA fibers, hot drawn 3 22. (o) Dried fibers, (Δ) fibers with 13% water (Garrett and Grubb, 1986, Figure 10a).

the fibers increased the storage modulus at RT by about 30% and also increased the glass transition by about 35
C. The morphology of the PVA fibers starts to develop during the spinning process and is further refined during the drawing. In some cases, such as wet or gel spinning, very little orientation and crystallization are introduced during the fibers formation, and most of the structure is developed during drawing process (Fig. 16.4; Yamaura and Kumakura, 2000). On the other hand, when water-plasticized melt spinning is performed, the process is performed under stress, and the dried fibers exhibit partial crystallinity. It can be seen in Fig. 16.5 that during wet spinning of water-plasticized PVA, only part of the final crystallinity was developed during the spinning stage and that additional crystallinity was developed and perfected during the hot drawing process, which also increased the melting temperature. The crystallization of the PVA in melt spun fibers, during the hot drawing process, is affected by several parameters. It has been demonstrated (Wu et al., 2010) that the main factors are the temperature, time, Dr, and draw rate. The approximate crystallinity of the PVA fibers was calculated from DSC curves by dividing the area of the melting peak by the melting enthalpy of 100% PVA crystals. It was found that when the fibers were heat treated at different temperatures without introducing stress by drawing, the crystallinity did not increase at 120
C

278

Thermal Analysis of Textiles and Fibers

Figure 16.4 DSC traces of undrawn and drawn PVA fibers, which were produced by gel spinning process (Yamaura and Kumakura, 2000, Figure 4). PVA88 is partially hydrolyzed grade, PVA99 is fully hydrolyzed grade, and PVA8899 is hydrolyzed during the drawing stage.

(393K). Once the temperature was increased to 150
C (423K) and to 180
C (453K), the crystallinity did increase over time, and these curves were marked “temperature-induced crystallinity” in Fig. 16.6 (Wu et al., 2010). When the fibers were drawn at different rates, the crystallinity kept increasing over time to higher values than the “temperature-induced crystallinity,” and these curves were marked as “stress-induced crystallinity.” It can be seen at all three temperatures that the crystallization increased faster as the drawing rate was increased. The crystallinity values in Table 16.1, which were calculated from the values in Fig. 16.6 (Wu et al., 2010), demonstrate the relative contribution of temperature, time, and stress. It was noted that at low shear rate of 50 mm/min the contribution of the stress-induced crystallization decreased with temperature, because the relaxation rate also went up at higher temperatures and offset the orientation effects of the drawing. Once the shear rate was increased to 100 mm/min, the stress contribution increased with temperature, because this effect became more pronounced than the molecular relaxation. As the shear rate was increased to 500 mm/min, the rate of the stress-induced crystallinity was the fastest, but the total developed

Thermal analysis of poly(vinyl alcohol) fibers

279

Figure 16.5 DSC runs of water-plasticized, melt spun PVA fibers (KURARAY POVAL 2899) before and after hot drawing. Dr 5 3.0 at the spinning stage was calculated as the ration of the take-up speed to the calculated jet speed. Dr 5 4.0 during the hot drawing at 200
C was the ratio of the fiber speeds before and after the drawing. Source: Ophir, Z., Jaffe, M., 2014. Report to United Soybean Board, Contract Project 14405125286.

crystallinity was less than was obtained at 100 mm/min, because the drawing time was obviously lower by the time the maximum Dr was reached. Additional information about the effect of drawing on melt spun PVA fibers can be seen in Fig. 16.7, which depicts thermomechanical analysis (TMA) traces of water-plasticized melt spun PVA fibers, after drying, before and after hot drawing. Both samples were loaded at the same stress level, and dramatic differences can be seen in the thermomechanical properties of the fibers. Both fibers deform very little below the glass transition temperature, but the difference between them becomes very clear above Tg. Once Tg is exceeded, the undrawn fiber becomes very soft and starts creeping, while the drawn fiber remains quite stiff up to 180
C and even shrinks somewhat, because of the relaxation in the amorphous regions, despite the tensile stress, which is imposed by the TMA Probe. DSC traces of the same fibers in Fig. 16.5 show that both the heat of fusion and melting temperature increased significantly during the hot drawing. However, the impact of these changes on the mechanical behavior of these fibers in actual use is not obvious.

280

Thermal Analysis of Textiles and Fibers

Figure 16.6 Crystallinity of PVA fibers (Wu et al., 2010, Figure 4), which were drawn at different temperatures and drawing rates. The approximate crystallinity values were calculated as the ratio between the DSC endothermic peak and the melting enthalpy of 100% crystalline PVA.

Table 16.1 Crystallinity values developed during hot drawing on top of the already existing crystallinity of the melt spun fibers. 120
C Stress (%) 50 mm/min (6 min) 100 mm/min (3 min) 500 mm/min (0.5 min)

Thermal (%)

150
C Total (%)

Stress (%)

Thermal (%)

180
C Total (%)

Stress (%)

Thermal (%)

Total (%)

8.7

0

8.7

8.3

3.3

11.6

1.7

5.0

6.7

10.3

0

10.3

12.5

1.7

14.2

18.0

2.9

20.9

9.0

0

9.0

9.2

0.3

9.5

9.2

0.5

9.7

Thermal-induced crystallinity developed at each temperature as a function of time without drawing, additional stress-induced crystallinity developed as a function of the drawing rate, and the total crystallinity is the sum of both as a function of temperature, time, and stress (Wu et al., 2010).

282

Thermal Analysis of Textiles and Fibers

Figure 16.7 TMA of water-plasticized, melt spun, and dried PVA fibers, before and after hot drawing. Source: Ophir, Z., Jaffe, M., 2014. Report to United Soybean Board, Contract Project 14405125286.

Dynamic mechanical analysis (DMA) of water-plasticized melt spun fibers (Fig. 16.8) and cast films (Fig. 16.9) show that the glass transition is spread over a wide range of temperatures and that center of the softening range in about 90
C, although the peak in tan δ is significantly higher. Some of the transition temperature spread can be attributed to the strong hygroscopic nature of PVA and evaporation of residual water during the test and also to the fact that industrial grade PVA is not a pure material and contains residual branching and acetate groups, which were not completely hydrolyzed. Shifts in the DMA spectra of water-plasticized melt spun fibers occur once the fibers are hot drawn (Fig. 16.10), and additional changes are also attributed to the heat treatment temperature and type of PVA (Fig. 16.11). The similarity between the PVA and PE molecules was seen as a potential for producing high-strength PVA fibers by gel spinning followed by very high degree of hot drawing. The advantages of PVA over PE molecules are both the higher melting point (i.e., higher use temperature) and better adhesion because of the abundance of hydrogen bonds (Kwon et al., 1984). As discussed earlier, attempts to achieve very high Drs, similar to those obtainable in PE, were only partially successful even when the drawing was performed very close to the melting point. Garrett and Grubb (1988) performed a series of dynamic mechanical

DMTA Temp Ramp PVA 28-99 melt spun fiber 197-17-4 10

10

10

0

9

–1

10

tan_delta ( ) ()

E′ ( ) (Pa)

10

8

10

10

7

0.0

20.0

40.0

60.0

80.0

100.0

120.0

140.0

160.0

10 180.0

–2

Temperature (ºC)

Figure 16.8 DMA of water-plasticized melt spun PVA fibers. Source: Ophir, Z., Jaffe, M., 2014. Report to United Soybean Board, Contract Project 14405125286. DMTA Temp Ramp PVA 28-99 film 10

10

10

0

9

–1

10

tan_delta ( ) ()

E′ ( ) (Pa)

10

8

10

10

7

0.0

20.0

40.0

60.0

80.0

100.0

120.0

140.0

160.0

10 180.0

–2

Temperature (ºC)

Figure 16.9 DMA of cast PVA film from water solution. Source: Ophir, Z., Jaffe, M., 2014. Report to United Soybean Board, Contract Project 14405125286.

284

Thermal Analysis of Textiles and Fibers

Dynamic Temp Ramp PVA 28-99 drawn fiber 197-18-5 10

10

9

E′ ( ) (Pa)

10

8

10

10

7

20.0

0.0

40.0

60.0

80.0

100.0

120.0

140.0

160.0

180.0

Temperature (°C)

Figure 16.10 DMA of water-plasticized melt spun PVA fiber after hot drawing 4 3 at 200
C. Source: Ophir, Z., Jaffe, M., 2014. Report to United Soybean Board, Contract Project 14405125286.

E′, E″ (dyne/cm2)

1011

1010

1

S–PVA

2

A–PVA

3

I–PVA

1 3

E′

2

3 2 1

109 E″

108 –100

–50

0

100 50 Temperature (°C)

150

200

Figure 16.11 DMA of heat-treated PVA films as a function of the PVA tacticity and heat treatment temperature (Nagai and Takayanagi, 1965). (1) S-PVA 207
C, (2) A-PVA 210
C, and (3) I-PVA 150
C.

thermal analyses (DMTAs) tests on fibers that were drawn at different Dr from 2 3 to 38 3 (Figs. 16.12 and 16.13). It was observed that beside some shifts of the glass transition, a second relaxation peak below the melting point not only moved to higher temperatures but also decreased in size as the Dr was increased.

Thermal analysis of poly(vinyl alcohol) fibers

285

12.0

λ = 38×

11.5

λ = 15×

Log(E′) (dyne/cm2)

11.0 λ = 8.5× 10.5 λ = 2× 10.0

9.5

9.0

8.5 –150 –100 –50

0

50

100

150

200

250

Temperature (ºC)

Figure 16.12 Storage modulus (E0 ) versus temperature spectra at 3.5 Hz for drawn PVA fibers at different draw ratios λ (Garrett and Grubb, 1988, Figure 5).

λ = 2×

λ = 5×

tan δ

0.2 λ = 8.5×

0.0 λ = 15×

λ = 38×

–150 –100 –50

0 50 100 150 Temperature (ºC)

200

250

Figure 16.13 Loss factor (tan δ) versus temperature spectra at 3.5 Hz for drawn PVA fibers at different draw ratios λ (Garrett and Grubb, 1988, Figure 4).

286

Thermal Analysis of Textiles and Fibers

0.05

0.04

tan δ

0.03

0.02

0.01

0

50

0

100

150

200

Temperature (ºC)

Endo

Exo

Figure 16.14 Temperature dependence of tan δ for the first, second, and third zone drawn fibers, and the three times zone drawn and vibrationally heat-treated fiber (order of curves from top to bottom) (Kunugi et al., 1990, Figure 12).

30

60

90

120

150

180

210

240

270

Temperature (ºC)

Figure 16.15 DSC curve for PVA fiber, which was zone drawn at 260
C. The fiber was finely cut and packed in the Al pan. Heating rate 10
C/min (Kunugi et al., 1990, Figure 9).

It was suggested by Garrett and Grubb (1988) that this loss of molecular mobility and its shifting to higher temperatures, once the Dr is increased, are the reasons for the inability to draw the fibers beyond certain limits. Garrett attempted to correlate these changes in the relaxation peak with other material factors, which

Thermal analysis of poly(vinyl alcohol) fibers

287

Figure 16.16 TGA of PVA films I—virgin (cast and dried), II—heat treated, and III—crosslinked PVA films (Gohil et al., 2005, Figure 10). (A) covers the low temperature range from 25-245C, (B) covers the high temperature range from 250-500C.

may cause them, such as crystallinity, orientation, melting temperature, and crystal thickness, but could not identify the actual cause for the loss of the molecular mobility. Kunugi et al. (1990) performed several steps of zone drawing, followed by heat treatment under vibration, in an attempt to maximize the Dr and the fibers mechanical properties. It was observed that the effects of these steps was to reduce the intensity of the glass transition relaxation without much change in this transition temperature, while the relaxation near the melting point both shifted to higher temperatures and decreased in size (Fig. 16.14). DSC traces of the highly drawn fibers showed a very sharp melting peak (Fig. 16.15). Kunugi suggested that the difference between the drawing behavior of PVA fibers and gel spun PE fibers is caused

288

Thermal Analysis of Textiles and Fibers

both by the existence of OH groups and the hardness of the crystals. OH groups cause steric hindrance and the hydrogen bonds create intra- and intermolecular interactions. Both mechanisms prevent slippages of molecular chains during drawing. The reduction of the size of the relaxation peak around the glass transition, after the fibers are drawn, is explained as inhibition of molecular movement of the amorphous chains. The changes in the relaxation peak in the vicinity of the melting

Figure 16.17 DSC curves of a—uncross-linked and b—cross-linked PVA films (Li et al., 2005, Figure 5).

Figure 16.18 Weight loss of uncross-linked and cross-linked PVA films at 100
C (Li et al., 2005, Figure 6).

Thermal analysis of poly(vinyl alcohol) fibers

289

point, together with the sharpness of the melting peak (Fig. 16.15), are explained by perfecting the crystallites and their uniformity. The production of PVA fibers ends in some cases by chemical cross-linking, which improves the fibers environmental stability. Fig. 16.16 (Gohil et al., 2005) depicts TGA curves of PVA films in various states: virgin (as-cast), heat treated, and cross-linked. It can be seen in these curves that heat treatment is more effective than chemical cross-linking in limiting the thermal decomposition at regular use temperatures below 150
C. However, chemical cross-linking still contributes to the stability of PVA fibers and films at regular use temperatures by reducing water pickup (Figs. 16.17 and 16.18; Li et al., 2005) and therefore improving the performance of PVA fibers in hot water.

16.4

Conclusion

Thermal analysis methods are widely used in characterization of PVA fibers throughout all stages of the production process from raw materials to the final products. It was found that each of these methods is most sensitive to different chemical and physical properties of the PVA fibers, and the use of combination of these methods provides a broader picture, than each method alone, of the processes and properties of PVA fibers. TGA is mostly used in assessing water pickup and thermal degradation of both the raw PVA and the final PVA fibers. DSC has found uses in characterizing materials and in optimizing processing conditions of all stages. TMA was found useful in characterization of creep properties of fibers, while DMA is mostly used in analysis of the effects of temperature through the use temperature spectrum and in better insight into the limiting factors and effects of hot drawing.

References Cebe, P., Grubb, D., 1985. J. Mater. Sci. 20, 4465. Cha, W.I., Hyon, S.H., Ikada, Y., 1994. J. Polym. Phys. 32, 297. Chen, N., Li, L., Wang, Q., 2007. Plast. Rubber Compos. 36 (7/8), 283. Garrett, P.D., Grubb, D.T., 1986. J. Mater. Res. 1 (6), 861. Garrett, P.D., Grubb, D.T., 1988. J. Polym. Sci.: B: Polym. Phys. 2509. Gilman, J.W., VanderHart, D.L., Kashiwagi, T., 1994. Chap. 11 in Fire and polymers II. In: ACS Symposium Series, vol. 599, p. 161. Gohil, J.M., Bhttacharya, A., Ray, P., 2005. J. Polym. Res. 13, 161. Jang, J., Lee, D.K., 2003. Polymer 44, 8139. Khasbiullin, R.R., Kostina, Y.V., Petrova, T.F., Bondarenko, G.N., Chalykh, A.E., Chuvaev, V.F., et al., 2014. Polym. Sci., Ser. A 56 (5), 569. Kunugi, T., Kawasumi, T., Ito, T., 1990. J. Appl. Polym. Sci. 40, 2101. Kwon, Y.D., Kavesh, S., Preversek, D.C., 1984. United States Patent 4,440,711. Li, L., Chen, N., Wang, Q., 2010a. J. Polym. Sci. B 48, 1946.

290

Thermal Analysis of Textiles and Fibers

Li, L., Chen, N., Wang, Q., 2010b. J. Polym. Sci.: B Polym. Phys. 48, 1946. Li, L., Wang, Q., Wang, R., 2005. J. Appl. Polym. Sci. 98, 774. Magalhaes, S.M., Toledo Filho, R.D., Rego Fairbrain, E.M., 2013. Rev. Mater. 18 (4), 1587. Mitchenko, Y.I., D’yachkov, A.N., Rudneva, L.D., 2011. Fiber Chem. 44 (1), 41. Nagai, A., Takayanagi, M., 1965. Kogio Kagaku Zasshi 68, 836. Ophir, Z., Jaffe, M., 2014. Report to United Soybean Board, Contract Project 14405125286. Peppas, N.A., Merrill, E.W., 1976. J. Appl. Polym. Sci. 20, 1457. Sakurada, I., 1985. Poly Vinyl Alcohol Fibers. Marcel Dekker Inc. Sakurada, I., Ito, T., Nakamae, K., 1967. J. Polym. Sci. Polym. Symp. 15, 75. Shaikh, R.P., Kummar, P., Choonara, Y.E., Du Toit, L.C., 2012. Biofabrication 4, 1. Shie, J.L., Chen, Y.H., Chang, C.Y., Lin, J.P., Lee, D.J., Wu, C.H., 2002. Energy Fuels 16, 109. Taghizadeh, M.T., Yeganeh, N., Rezaei, M., 2015. J. Appl. Polym. Sci. 42117, 1. Tashiro, K., Kobayashi, M., Tadokoro, H., 1977. Macromolecules 10 (4), 731. Uddin, A.J., Narusawa, T., Gotch, Y., 2011. Polym. Eng. Sci. 51 (4), 647. Wu, Q., Chen, N., Wang, Q., 2010. J. Polym. Res. 17, 903. Wu, W., Tian, H., Xiang, A., 2012. J. Polym. Environ. 20, 63. Yamaura, K., Kumakura, R., 2000. J. Appl. Polym. Sci. 77, 2872. Yumin, X., Zhixin, W., Huigong, T., Yimin, W., 2009. e-Polymers No. 001.

Further reading Chalykh, A.E., Chuvaev, V.F., Gerasimov, V.K., 2014. Polym. Sci. 56 (5), 569.

Polybenzimidazole fiber

17

Joseph D. Menczel Thermal Measurements LLC, Fort Worth, TX, United States

Abstract Polybenzimidazole (PBI) [poly[2,20 -(m-phenylene)-5,50 -bis-benzimidazole]] is a high-performance heterocyclic polymer fiber for use at very high temperature. It is made from tetraaminobiphenyl-(3,30 -diaminobenzidine) and diphenyl iso-phthalate. It is stable at high temperatures and has excellent resistance against aggressive chemicals. It is not flammable and does not readily ignite. The limiting oxygen index of PBI (the minimum oxygen concentration required to sustain steady-state burning) is 58%59%, so it will not burn in air. The short-term thermo-oxidative stability of PBI is the highest among all plastics. Its radiation resistance is also good. PBI is used in various different physical forms, such as compression molded resin pieces, composites, fiber, and films. It is also used to manufacture protective apparel, such as astronaut space suits, firefighter’s gear, high-temperature protective gloves, aircraft wall fabrics, and welders’ apparel. Lately, it is used as a membrane in high-temperature fuel cells. The lifeline that connects the astronauts to the spaceship is made of PBI. Obviously, the disadvantage of PBI is its high price.

Polybenzimidazole (PBI) [poly[2,20 -(m-phenylene)-5,50 -bis-benzimidazole]] is a high-performance heterocyclic polymer fiber for use at very high temperatures. It is made from tetraaminobiphenyl-(3,30 -diaminobenzidine) and diphenyl iso-phthalate. It is stable at high temperatures and has excellent resistance against aggressive chemicals. It is not flammable and does not readily ignite. The limiting oxygen index of PBI (the minimum oxygen concentration required to sustain steady-state burning) is 58%59%, so it will not burn in air. The short-term thermo-oxidative stability of PBI is the highest among all plastics (see Table 17.1). Its radiation resistance is also good. PBI is used in various different physical forms, such as compression molded resin pieces, composites, fiber, and films. It is used to manufacture Thermal Analysis of Textiles and Fibers. DOI: https://doi.org/10.1016/B978-0-08-100572-9.00017-3 © 2020 Elsevier Ltd. All rights reserved.

292

Thermal Analysis of Textiles and Fibers

Table 17.1 Short-term thermo-oxidative properties of polybenzimidazole fiber. Temperature ( C)

Time (useful life)

600 450 400 330 300

35 s 5 min 1h 24 h 1 week

Source: From Sandor, R.B., 1990. High Perform. Polym. 2, 25: Table 2 with permission of High Performance Polymers.

Figure 17.1 DSC curves of as-spun and drawn PBI fiber in the free-to-shrink and constrained (fixed-length) modes. Tg of the drawn fiber in constrained mode is the highest, Tg of the as-spun fiber is the lowest. DSC curves were recorded on a TA Instruments 2100/ 910 DSC at a heating rate of 20 C/min. Source: Springer from Menczel, J.D., 2000. J. Therm. Anal. Calorim. 59, 1023. Reprinted with permission of Springer.

protective apparel, such as astronaut space suits, firefighter’s gear, high-temperature protective gloves, aircraft wall fabrics, and welders’ apparel. Lately, it is used as a membrane in high-temperature fuel cells (Cassidy, 1980; Levine, 1969; Powers and Serad, 1986). The lifeline that connects the astronauts to the spaceship is made of PBI. Obviously, the disadvantage of PBI is its high price. PBI is not white like most plastics are; PBI pieces or fibers usually have yellow or brown color. It retains mechanical integrity up to 400 C. PBI can be readily dissolved in strong protonic acids (like sulfuric acid or methanesulfonic acid), but its solubility in weaker acids is controversial: several authors reported complete solubility in formic acid, others observed partial solubility only.

Polybenzimidazole fiber

293

PBI has high moisture regain (B15%) that makes the clothing made of PBI comfortable to wear (the moisture regain of cotton is around 16%, while that of nylon can go up to B8%). Most frequently, PBI is used as a fiber. PBI is dry spun from a 25-mass% dimethylacetamide solutions containing lithium chloride (to reduce the dope viscosity, Conciatori et al., 1967), although the technology exists also for wet spinning (Santangelo, 1964). After the solvent is evaporated the fiber is washed with water to remove the residual solvent, phenol, and LiCl; then it is drawn at elevated temperatures to improve the mechanical properties. Finally, it is sulfonated to reduce its flame shrinkage (Sandor, 1990), and most often made into staple. The first application of PBI was fire blocking and thermal-protective apparel. After the 1990s the application spread to solid physical parts and membranes. PBI will not flow and therefore, for bulk pieces, it needs to be compression molded. Molded PBI plastic pieces have high strength and low weight. Especially the compression strength and modulus are high: it has the highest compressive strength of any unfilled plastic. PBI molded pieces can be used up to 300 C350 C. In addition, these molded parts have high dimensional stability and good retention of electrical properties at high temperatures. The notch sensitivity of PBI is considerable, so all the surfaces of PBI parts must be smooth and corners should be radiused. The bulk PBI pieces are very hard, and diamond tools are needed for processing. Molded PBI has a trade name of Celazole. Commercial PBI is amorphous, and its glass transition temperature is the highest among commercial polymers. When recording the differential scanning calorimetry (DSC) curve of PBI, one gets a very broad water evaporation peak between room temperature and 250 C. Then, as shown in Fig. 17.1, the glass transition follows. Tg of the as-spun fiber is 387 C. Tg of the drawn fiber is higher, and it depends on the mode of the measurement: in free-to-shrink mode, Tg is 401 C, while in the constrained mode, it is 435 C (Menczel, 2000). This was the first time when Tg of an amorphous drawn fiber recorded in the free-to-shrink state could be compared to Tg of the fiber measured at fixed length. Like the melting point, the glass transition temperature also depends on possible shrinkage of the fiber during the measurement. The reason is clearly the difference in the change of the free volume for the mentioned two cases. Dynamic mechanical analysis (DMA) curves of the as-spun and drawn fibers are shown in Figs. 17.2 and 17.3. The primary relaxation (the glass transition) of the as-spun fiber occurs at 465 C, higher than the DSC Tg due to the frequency effect. As can be seen in Figs. 17.2B and 17.3, three sub-Tg relaxations were recorded for both the as-spun and the drawn fiber, although with different intensities. The lowest temperature relaxation at 294 C may correspond to onset of rotation of the m-phenylene ring. It is more difficult to assign the remaining two relaxations. The β-relaxation has smaller intensity in the drawn fiber, and it is split. This may indicate that this relaxation can be the consequence of water evaporation, and more than one state of bond water may exist in PBI. At the same time, the relaxation at the 6 C could be assigned to onset of rotation of the benzimidazole ring and melting of free water.

294

Thermal Analysis of Textiles and Fibers

Figure 17.2 (A) DMA curves of as-spun PBI fiber (E0 is the Young’s modulus, tan δ 5 Ev/E0 ). The first drop in the tensile modulus at around 0 C is due to the high moisture content of the fiber. The second drop in E0 at around 250 C is due to release of bound water. Tg is around 465 C (higher than in the DSC measurements due to the frequency effect). Measurements carried out on a Rheometrics RSA2 Solid Analyzer at 1 Hz. (B) The tan δ versus temperature curve of as-spun PBI fiber at high sensitivity. The lowest temperature peak is the onset of rotation of the m-phenylene ring, the peak at 6 C may correspond to start of rotation of the benzimidazole ring/melting of free water, and the 240 C peak somehow reflects the release of bound water from PBI. Measurements carried out on a Rheometrics RSA2 Solid Analyzer at 1 Hz. Source: Springer from Menczel, J.D., 2000. J. Therm. Anal. Calorim. 59, 1023. Reprinted with permission of Springer.

Polybenzimidazole fiber

295

Figure 17.3 DMA curves of drawn PBI fiber (E0 and tan δ). Due to the frequency effect, Tg is so high that the fiber broke at ca. 450 C. The assignment of the tan δ peaks is the same as in Fig. 17.2B, but the water release peak splits into two peaks. This may indicate that the release of bound water may be different in the fiber regions with different orientation. Measurements carried out on a Rheometrics RSA2 Solid Analyzer at 1 Hz. Source: Springer from Menczel, J.D., 2000. J. Therm. Anal. Calorim. 59, 1023. Reprinted with permission of Springer.

References Cassidy, P.E., 1980. Thermally Stable Polymers. Marcel Dekker, New York. Conciatori, A.B., Chenevey, E.C., Bohrer, T.C., Prince, A.E., 1967. J. Polym. Sci., C 19, 49. Levine, H.H., 1969. Encycl. Polym. Sci. Technol. 11, 188. Menczel, J.D., 2000. J. Therm. Anal. Calorim. 59, 1023. Powers, E.J., Serad, G.A.,1986. History and Development of Polybenzimidazoles. In: Seymour, R.B., Kirshenbaum, G.S. (Ed.), High Performance Polymers, Their Origin and Development, Proc. Symp. History of High Perf. Polym. at ACS Meeting, 1518 April 1986, Springer, New York. pp. 355373 Sandor, R.B., 1990. High Perform. Polym. 2, 25. Santangelo, J.G., 1964 (Celanese Corp.) Process for Wet Spinning Polybenzimidazoles, U.S. Patent 3,441,640. Hoechst Celanese. Properties of polybenzimidazole high-performance fiber. Tech. Bull.

Further reading Reinhart, K.A., Powers, E.J., Calundann, G.W., Driscoll, C.P., January 1974. Technical Report ASD-TR-49. Hoechst Celanese. Celazole. Tech. Bull.

Thermal analysis of acrylic and carbon fibers

18

Shahram Arbab1 and Joseph D. Menczel2 1 Textile Engineering Department, ATMT Research Institute, Amirkabir University of Technology, Tehran, Iran, 2Thermal Measurements LLC, Fort Worth, TX, United States

Abstract In this chapter, first, we describe how the various thermal analysis techniques can be used to characterize the PAN homopolymer or copolymer fibers, and then we show how these techniques can be applied to study the transformation of PAN to CFs, because the CF production can be considered, and is a kind of thermal degradation of PAN. As usual, differential scanning calorimetry (DSC), thermogravimetric analysis (TGA), thermomechanical analysis (TMA), and dynamic mechanical analysis (DMA) will be used in both areas.

18.1

Overview

Polyacrylonitrile (PAN) fiber is one of the most important textile fibers. In addition, it is widely used as a precursor for manufacturing high-performance carbon fibers (CFs). A significant difference between these two applications is that the source of the CF production is often PAN homopolymer, but in the textile industry, acrylonitrile (AN) copolymers are used. In this industry, two kinds of acrylic fibers are popular: the acrylics with the AN content of at least 85%, and modacrylics that contain 35%
85% AN, and the remaining 15%
65% is usually composed of halogenated monomers to make the fibers flame retardant. For CFs the mechanical properties play a significant role. In this field the thermal analysis research studies concentrated mostly on control of the rate and the heat of the exothermic reactions taking place during the stabilization. Despite numerous efforts, there is still no clear approach on optimization of the stabilization process. The main reasons for this are the following: 1. complexity of the stabilization reactions, 2. effect of chemical composition on the progress of chemical reactions leading to CFs and complicated structure of the precursor fiber, and 3. difficulty of separating the chemical reactions taking place during the “cyclization” step, since a number of such reactions take place simultaneously.

In this chapter, first we describe how the various thermal analysis techniques can be used to characterize the PAN homopolymer or copolymer fibers, and then we show how these techniques can be applied to study the transformation of PAN to CFs, because the CF production can be considered and is a kind of thermal degradation of PAN. As usual, differential scanning calorimetry (DSC), thermogravimetric Thermal Analysis of Textiles and Fibers. DOI: https://doi.org/10.1016/B978-0-08-100572-9.00018-5 © 2020 Elsevier Ltd. All rights reserved.

298

Thermal Analysis of Textiles and Fibers

analysis (TGA), thermomechanical analysis (TMA), and dynamic mechanical analysis (DMA) will be used in both areas.

18.2

Introduction

PAN fiber is the third major textile fiber after polyesters and nylons. The average molecular mass of PAN for fibers is around 100,000. It is also widely used as a precursor for manufacturing advanced materials, such as high-performance CFs (Bacon, 1960; Bashir, 1991; He, 2005; Morgan, 2005; Ouyang, 2008; Shen, 2008; Frank et al., 2012). The first application of PAN fiber happened in the textile industry, but later PAN also became the main precursor for high strength CFs (Paiva et al., 2003): the properties of CFs are strongly affected by the structure and thermal properties of the precursor fibers (Mittal et al., 1998; Thu¨nemann and Ruland, 2000). Therefore detailed thermal characterization of PAN fibers is crucial for both, manufacturing textile fibers with the desired properties and designing the CF production line. The worldwide global acrylic and modacrylic fiber production in 2014 was 1.8 million metric tons, this was around 3% of all the synthetic fiber production (of course, poly(ethylene terephthalate) (PET) was the leader with 77%). Some of the well-known commercial names of acrylic fibers are Cortelle, Acrilan, Zefran, Sayelle, Orlon, Dralon, and Drytex, while common names of modacrylic fibers include Teklan, Dynel, SEF, Zefran, Verel, and Elura. Twenty-five percent of these fibers are monofilament and 75% are staple. The vast majority of these fibers are manufactured in Asia. As mentioned earlier, in acrylic fibers, the AN content is at least 85%, and the remaining 15% is vinyl acetate (VAc) or methyl acrylate to increase solubility of PAN for the spinning or improve dyeing properties of the fibers. In the modacrylic fibers the AN content is between 35% and 85%, and the fiber contains 15%
65% halogenated comonomers to make the fibers flame retardant (Frushour and Knorr, 1985). Most of these fibers are used for apparel, so the tensile properties are low, their tenacity is between 2.7 and 3.6 cN/ dtex. PAN is moisture sensitive. Its flame retardancy is good, because it usually contains halogenated comonomers, as mentioned earlier. The light resistance of acrylic fibers is also good, so is their resistance to chemicals. The majority of CFs on the market are made from PAN, but not all. About 10% of all CF is manufactured from rayon and petroleum pitch; but special CFs are also made from other polymers, such as polyimides and polybenzimidazole. Vaporgrown carbon nanofibers (carbon nanotubes) are the latest addition to the family of CFs. All these CFs are much more expensive than the CF made of PAN. PAN fibers need to be thermally stabilized by heating them in an oxygencontaining atmosphere at temperatures of 220 C
330 C for some time. This way an oxidized PAN fiber is formed, that later can be transformed into CF at temperatures above 1000 C by controlled pyrolysis. CF is a fiber containing at least 90% carbon. Graphite fibers have carbon content over 99%. CFs are used in military and civil aircrafts. About 25% of the body of civil aircrafts are made of CFs. In addition, CFs are popular in high-performance composites (aerospace, military,

Thermal analysis of acrylic and carbon fibers

299

Table 18.1 Comparison of mechanical properties of various synthetic fibers. Fiber

Tensile modulus (GPa)

Tensile strength (GPa)

Elongation (%)

Aramid PET Nylons Gel spun polyethylene (e.g., Spectra) CF from PAN CF from mesophase pitch

65 20 5 7

3 1 0.15
1 3

5 20
150 25 3

400 500

4 3

0.5 1

automobiles, motorsports, special bicycles, various fiber
reinforced plastics, etc.). However, CF is not cheap that limits its application. The chemical resistance of CFs is excellent, they have low density and low thermal expansion, and good tolerance to high temperatures. Their stiffness and tensile strength are very high. A microstructure peculiarity of CFs is that the carbon atoms are bonded in various crystalline cells. These cells are arranged parallel to the long axis of the CF, and this is the major reason for the high strength/volume ratio of these fibers. The CF “yarn” is called “a tow,” and it consists of several thousand carbon single fibers. The diameter of the carbon monofilaments is around 5 μm. The tows may be woven to form a fabric. Drawn PAN fiber is used for CF production, because the higher tensile strength precursor results in the higher tenacity CF. Carbonization under strain significantly improves tensile strength and modulus of CFs. The tensile strength and the Young’s modulus of CFs are influenced by the type of precursor and the temperature of heat treatment. CFs are very brittle. Since CFs have sheet-like structure, crack propagation is very easy in them. Some mechanical properties of CFs are compared to other fibers in Table 18.1. CFs consist of hexagonal sheets of carbon atoms, such as graphite. CFs may be turbostratic (PAN type), graphitic (from mesophase pitch), or the structure may be hybrid consisting of both mentioned types. In CFs of turbostratic structure the sheets of carbon atoms are irregularly folded. The CFs of turbostratic structure have higher tensile strength, but the CFs made from mesophase pitch have higher thermal conductivity and higher Young’s modulus. The structure of CFs is usually characterized by X-ray scattering, electron microscopy, and atomic force microscopy (AFM). Strictly speaking, there is no three-dimensional order in most CFs, the structure is a kind-of “mesophase” structure (similar to discotic liquid crystals). In PAN-based CFs, basal planes are formed (these planes are perpendicular to the principal axis). The basal plane orientation will increase with the increase of the carbonization temperature. It is also important that the orientation in the CFs is not homogeneous: the skin has higher axial orientation than the core. The production volume of CFs in 2012 was 112,000 metric tonnes (MT), and 170,000 MT is expected for 2020.

300

Thermal Analysis of Textiles and Fibers

Low modulus and low-strength CFs are manufactured at low carbonization temperature, around 1000 C. Depending on the type of CF, the high-heat treatment (carbonization) temperature can be as high as 2000 C
3000 C [production of highmodulus (HM) fibers]. If the heat treatment (carbonization) temperature is between 1500 C and 2000 C, the highest tensile strength ( . 5.5 GPa) fibers are produced, while CF produced at 2500 C
3000 C (graphitization) will give fibers of the highest tensile modulus (higher than 450 GPa). Thus the CFs can be classified based on their tensile modulus. Ultra-HM-type (UHM) fibers have moduli higher than 450 GPa; the modulus of HM fibers is between 200 and 350 GPa. Low modulus fibers have a modulus ,100 GPa. The first step in the CF production is oxidative stabilization (cyclization) at 220 C
330 C. The result of this heat treatment is a cyclic polymer that is nonplastic (it does not melt anymore). The second step is the carbonization (sometimes graphitization) at temperatures between 1000 C and 3000 C depending on the mechanical properties of the produced fiber. Most of the noncarbon elements are removed in the carbonization step, and several hours are required for this. CF is very expensive. An automobile part made from CF may cost 10 times as much as a similar metal part, but the price of CF has a decreasing tendency (15 years ago, this 10 times ratio was higher than 30). But the advantage of using CFs in automobiles appears when the weight of the car is looked at. The use of next generation of carbon-fiber composites may reduce the weight of automobiles by about 50%. During the oxidative stabilization step, the linear structure of PAN changes to an intermediate cyclic structure, so this step is often called “cyclization.” The cyclic structure has a critical role in posterior processing steps, such as carbonization, graphitization, and hydrolysis (Bashir, 1991; He et al., 2005; Morgan, 2005; Shen et al., 2008; Ouyang et al., 2008a,b; Ouyang et al., 2009; Wangxi et al., 2003; Rangarajan et al., 2002; Kakida and Tashiro, 1997; Yusof and Ismail, 2012; Mirbaha et al., 2013). Thermal stabilization is a time and energy-consuming step in the production of CFs from PAN (Yusof and Ismail, 2012). During this step, PAN fibers undergo complicated chemical reactions, namely, cyclization, dehydrogenation, and oxidation (Sen et al., 2003; Rahaman et al., 2007; Arbab and Zeinolebadi, 2013). The progress of the stabilization reactions and the intermediate structures formed during this step strongly affects the final properties of the CFs (Frank et al., 2012; Arbab and Zeinolebadi, 2013, 2017; Huang, 2009; Ju et al., 2014). Therefore optimization of the stabilization process is important. It is not as important as the final carbonization heat treatment but plays a significant role for determining the properties of the final CF and the possible reduction of the high production cost (Fitzer, 1986). Warner et al. (1979) introduced a model to describe the structure of oriented PAN fibers. They suggested that all the areas in the PAN fibers form two regions: partially ordered (paracrystalline) and disordered (amorphous). In the paracrystalline regions the chain assumes a distorted helical shape, and the result is the

Thermal analysis of acrylic and carbon fibers

301

˚ . Inside these rods, the nitrile units are formation of rods with a diameter of B6.0 A somewhat oriented, but their position and direction are irregular on the surface of the rod. The nitrile groups on the surfaces of neighboring rods can interpenetrate, and they form dipole pairs. Since the rods are ordered into a nematic-type array, the nitrile groups form a lamellar-type structure. Some support for this model was found in small angle X-ray scattering (SAXS) and wide angle X-ray scattering (WAXS) experiments of heat-treated fibers.

18.3

Production of polyacrylonitrile fibers

The acrylic fibers are usually produced by solvent spinning, that is, wet spinning or dry spinning (Salem, 2001). The solvents are dimethylformamide, dimethylacetamide, or dimethyl sulfoxide. Depending on the application of the fiber, the spinning is followed by drawing and possibly by texturing. In wet spinning, as soon as the polymer solution is extruded into the coagulation bath, phase separation starts taking place due to solvent/nonsolvent exchange and polymer chains precipitate as porous fibrils. The structure of the porous fibrils is controlled by the rate and kinetics of phase separation and gelation (Salem, 2001; Chen et al., 2006; Wang et al., 2007; Masson, 1995). The as-spun fibers contain numerous macro- and nanovoids (Arbab et al., 2008). The total porosity and the size of the voids determine drawability and thermal properties of the fiber in the next steps (Arbab et al., 2008, 2011a,b).

18.4

Thermal analysis of polyacrylonitrile and polyacrylonitrile fibers

18.4.1 Thermal analysis of polyacrylonitrile Frushour and Knorr (1985) gave a detailed review of the chemistry of PAN fibers and included some information on the thermal properties. PAN degrades at tem  peratures below its equilibrium melting point (Tm ); therefore Tm had to be estimated through various extrapolations. Krigbaum and Tokita (1960) determined 317 C, Slade (1970) 322 C, Berger et al. (1973) 330 C
335 C, Dunn and Ennis (1970) 326 C, Hinrichsen 320 C 6 5 C, and Frushour 344 C. All these data lead to an average of 330 C 6 17 C. This is a fairly good number, but there are problems with the determination of the heat of fusion of 100% crystalline PAN (equilibrium heat of fusion). Most authors used the copolymer method to determine the equilibrium heat of fusion. The equilibrium heat of fusion obtained was 45
100 J/g. In the copolymer method, it is assumed that the comonomer entities are excluded from the crystals. The mentioned authors (Krigbaum and Tokita, 1960; Slade, 1970; Berger et al., 1973) used different comonomers (isobutylene, styrene, and vinyl acetate (VAc)), and the significant scattering of the data indicates that the comonomers may be partially incorporated into the crystals.

302

Thermal Analysis of Textiles and Fibers

The glass transition temperature (Tg) of PAN is another serious issue. Of course, it depends on the technique used, that is, DSC, TMA, or DMA. Most DSC and TMA measurements gave Tg values between 85 C and 95 C. Dilatometrically, the Tg was obtained as 104 C (Krigbaum and Tokita, 1960), but DMA and dielectric analysis (DEA) give higher Tgs due to the frequency effect. Bohn et al. (1961) plotted the X-ray D-spacing as a function of temperature and found a break in the slope of the curve at around 90 C. They also found a break in the slope of the coefficient of thermal expansion versus temperature dependence at 85 C. If these two breaks reflect the same phenomenon, this can be interpreted as a “pseudo-glass transition” of the laterally bonded crystalline structure. The question arises whether this phenomenon is similar to the mesophase in polypropylene. Fast scan DSC measurements of Schick et al. (2017) tend to indicate that the mesophase in polypropylene is formed because of the enormous homogeneous nucleation density. This density can be so high that the space for crystal growth is extremely limited, so the crystal growth is terminated at the start of the crystallization process, and with a large amount of the rigid amorphous fraction (Menczel and Wunderlich, 1981): the whole sample becomes a material with a vast rigid amorphous content. Andrews and Kimmel (1965) interpreted the PAN phase as a “doubly bonded single phase.” They called this phase as glassy with some degree of orientation and molecular packing. Thus none of the thermal transitions of PAN is clear. In addition to the melting parameters, the “glass transition” (the 110 C transition taking place between 85 C and 140 C) of PAN is probably a paracrystalline transition. PAN is often fully “paracrystalline.” The sample variations make this situation more difficult. The most probable explanation to the mentioned controversy is that the Tg of PAN is somewhere around 100 C, but it can be overlaid by a paracrystalline transition. So, two transitions can exist in the 100 C temperature region. Controversial information is only available about these transitions. Since the major application of PAN is apparel, it is important to understand how various comonomers and water affect the glass transition. As mentioned earlier, serious confusion exists in the literature about the 110 C and 160 C transition (see the “DMA characterization” section). Many DSC measurements suggest that the 85 C
110 C transition is “glass transition.” Howard (1961) made a series of measurements on Tg of AN sample with vinyl acetate (VAc) comonomers. Tg was constant between 0% and 27% comonomer content. At higher VAc content, the 1/Tg value changed linearly with the VAc content, and Howard determined 110 C as Tg of the amorphous PAN homopolymer using the Fox
Lashaek equation (Fox and Lashaek, 1955). Frushour had similar results (Frushour and Knorr, 1985). Tg of PAN depends on the water content, water plasticizes this polymer, Tg may go down as low as 36 C for an AN
VAc copolymer with 7% VAc content (Frushour and Knorr, 1985). However, other authors concluded on the basis of DMA studies that the 110 C transition is a transition in the paracrystalline phase, whereas Tg is at 160 C (see the “DMA characterization” section). On the other hand, modulated TMA measurements indicate the opposite (Menczel, 2020 to be published). Presently, it is impossible to decide on this question.

Thermal analysis of acrylic and carbon fibers

303

18.4.2 Thermal analysis of polyacrylonitrile fibers 18.4.2.1 Differential scanning calorimetry characterization One of the first DSC curves of PAN fibers is shown in Fig. 18.1 (Dunn and Ennis, 1971). They stated that the melting point of PAN can be measured if fast heating is used, and it was obtained at 326 C. Other DSC curves of PAN fibers recorded at different heating rates in air and nitrogen are shown in Fig. 18.2. Characteristic temperatures of the exothermic peaks are summarized in Table 18.2. All PAN fibers exhibit an exothermic peak in the range of 200 C
400 C. It is attributed to stabilization reactions (Ouyang et al., 2008a,b; Sa´nchez-Soto et al., 2001; Catta Preta et al., 2007; Sviridov et al., 2009). As it can be seen in Fig. 18.2 and Table 18.2, increasing heating rate in both air and nitrogen shifts the initiation, maximum, and final temperatures of the exothermic peaks to higher temperatures. When the DSC measurements are carried out at heating rates of 2 and 5 C/min in air, the exothermic peak of the PAN/IA/MA fiber is broad and double (a shoulder is overlaid on the major peak) (Fig. 18.2A). As the heating rate increases to 10 C/ min, the shoulder disappears and the peak becomes sharper and perhaps, narrower. Hence, controlling the stabilization reactions and obtaining proper structure is more

Figure 18.1 Comparison of DSC curves of commercial acrylic fibers: (A) Orlon, (B) Cashmillion, (C) Crelan 61B, (D) Acrilan 16. The measurements were done on a DuPont (presently TA Instruments) 900 DSC. Free-to-shrink measurements, sample size 2
3 mg. Measurements in air heating rate were 2 C/min, endo down. Source: From Dunn, P., Ennis, B.C., 1970. J. Appl. Polym. Sci. 14, 1795. Reprinted with permission from Wiley.

Figure 18.2 DSC curves of PAN fibers at different heating rates in air and nitrogen atmospheres. Free-to-shrink measurements. (A) PAN/IA/MA, (B) PAN/IA, and (C) PAN/ VAc (IA: itaconic acid, MA: methyl acrylate). Solid curves: 2 C/min heating rate; double dotted curves: 5 C/min heating rate: dashed curves: 10 C/min heating rate. TA Instruments DSC 2910. Source: From Arbab, S., Zeinolebadi, A., 2017. Polym. Degrad. Stab. 139, 107. Reprinted with permission from Elsevier. Table 18.2 Characteristic temperatures of differential scanning calorimetry curves of polyacrylonitrile fibers in air and nitrogen atmospheres at heating rates of 2, 5, and 10 C/min. Fibers

PAN/IA/ MA PAN/IA

PAN/VAc

Characreristic temperatures



Ti ( C) Tpk1 ( C) Tpk2 ( C) Ti ( C) Tpk1 ( C) Tpk2 ( C) Ti ( C) Tpk ( C)

Rate 5 2 C/min

Rate 5 5 C/min

Rate 5 10 C/ min

Air

Nitrogen

Air

Nitrogen

Air

Nitrogen

175 250 302 210 264 306 240 290

210 258
210 258
240 300

180 270 328 215 284 328 245 310

220 274
215 273
245 328

205 282
230 286 346 250 346

225 290
230 288
250 342

Ti, Initiation temperature; Tpk1, first peak temperature; Tpk2, second peak temperature. Source: From Arbab, S., Zeinolebadi, A., 2017. Polym. Degrad. Stab. 139, 107. Reprinted with permission of Elsevier.

Thermal analysis of acrylic and carbon fibers

305

difficult at higher heating rates. In nitrogen, the exothermic peak is still double, but the shoulder is much less intense. Since oxidation reactions do not occur in nitrogen atmosphere, this peak is mainly attributed to cyclization and dehydrogenation reactions as well as chain rupture, which is more intense at higher temperatures (Kakida and Tashiro, 1997). However, some parts of the dehydrogenation reactions are absent in nitrogen. These are the reactions that result from oxygen attack to chains, leading to hydrogen release in the form of H2O. This may be a reason for narrower exothermic peak in nitrogen compared to air. As the heating rate increases, the exothermic peak of PAN/IA fibers in both air and nitrogen shifts to higher temperatures (Fig. 18.2B, Table 18.2). However, the shoulder of exothermic peak in air can still be observed at heating rate of 10 C/ min, but the ratio of the shoulder’s height to first peak’s height at heating rate of 10 C/min is much lower than those at heating rates of 2 and 5 C/min. This reveals that the number of reactions occurring in this temperature range decreased at the heating rate of 10 C/min. Some researchers believe that oxidation reactions occur in this temperature range (Ouyang et al., 2008a,b; Wangxi et al., 2003; Liu et al., 2009). Oxidation reactions are mainly controlled by rate of diffusion of oxygen into the fibers. Therefore changing the heating rate has less effect on their progress. In addition, oxygen uptake is decreased at higher heating rates because of the formation of a dense skin. As a result, the required time for oxygen to penetrate into core of fibers increases (Asakawa et al., 2012). The DSC exothermic peak of PAN/VAc fibers at heating rates of 2, 5, and 10 C/ min is narrow and the shoulder is weak (Fig. 18.2C). As expected, the initiation and the peak temperatures shift to higher temperatures (Table 18.2). The increase in height of exothermic peaks of PAN/IA/MA and PAN/IA fibers at higher heating rates is more significant compared to PAN/VA fibers (Fig. 18.2). The difference between thermal behaviors of fibers can be a result of the difference in nature of comonomers and some variations in molecular structure, such as molecular mass, molecular mass distribution, and spinning conditions. Characteristics of PAN/IA/MA and PAN/IA fibers are somewhat analogous and both of them contain acidic comonomers. These factors result in similarity in their thermal behavior. However, PAN/VAc fibers do not contain acidic comonomers and their characteristics, such as molecular mass, intrinsic viscosity, and sonic modulus, are different from PAN/IA/MA and PAN/IA fibers. These differences lead to variations in thermal behavior of fibers, such as lower effect of heating rate on exothermic stabilization reactions of PAN/VAc fibers. In addition, Fig. 18.2 suggests that the second exothermic peak in DSC curves of PAN/IA/MA (Fig. 18.2A) and PAN/IA (Fig. 18.2B) fibers is related to the presence of acidic comonomers. Thermal behaviors of the three fibers in air are also compared in Fig. 18.3. It can be seen that the exothermic reactions in PAN/VAc start at a higher temperature (240 C) compared to PAN/IA/MA and PAN/IA (175 C and 210 C, respectively). The exothermic peak in PAN/VAc is sharp and single compared to broad double peaks of PAN/IA/MA and PAN/IA fibers. This shows that in PAN/VAc the chemical reactions occur almost simultaneously and liberated heat in unit time is higher than in the PAN/IA/MA and PAN/IA fibers. The differences in thermal behavior of

306

Thermal Analysis of Textiles and Fibers

Figure 18.3 DSC curves of fibers in air atmosphere, (A) PAN/IA/MA, (B) PAN/IA, and (C) PAN/VAc fibers. TA Instruments DSC 2910. Source: From Arbab, S., Zeinolebadi, A., 2013. Polym. Degrad. Stab. 98, 2537, Fig. 8. Reprinted with permission from Elsevier.

fibers can be attributed to the chemical composition of polymers. The effect of VAc comonomers on decreasing the initiation temperature of stabilization reactions in PAN/VAc is not as much as itaconic acid comonomers in PAN/IA/MA and PAN/IA fibers. This can be attributed to slow elimination of acetate anion, which is the rate-determining step for the initiation of stabilization reactions (Liu et al., 2009). The presence of two acidic groups in the structure of itaconic acid comonomers reduces the initiation temperature of reactions and intensity of heat release significantly.

TGA measurements with various heating rates in air and nitrogen TGA curves of PAN fibers at different heating rates in air and nitrogen are shown in Fig. 18.4. Mass loss of fibers at the characteristic temperatures of the DSC curves is given in Table 18.3. According to Fig. 18.4, the mass loss of fibers at heating rates of 2 and 5 C/min in air atmosphere is similar. The maximum difference between mass losses at these heating rates is 3%. Initial mass loss of fibers at a heating rate of 10 C/min is similar to heating rates of 2 and 5 C/min. But at a specific temperature (280 C for PAN/IA/MA, 275 C for PAN/IA, 315 C for PAN/ VA), the slope of the mass loss curve increases sharply leading to significant mass loss compared to heating rates of 2 and 5 C/min. Therefore it may be suggested that at a rate of heating of 10 C/min, certain reactions leading to sudden mass losses (all types of dehydrogenation and chain rupture) are enhanced at or above the mentioned temperatures. Regarding the DSC curves shown in Fig. 18.2, increasing heating rates lead to more concurrent reactions. Polymer chains undergo more damage and chain rupture at higher heating rates; therefore higher mass losses are observed. In addition, at higher heating rates, there is an upward shift in the

Figure 18.4 TGA curves of PAN fibers at different heating rates in air and nitrogen. (A) PAN/IA/MA; (B) PAN/IA; (C) PAN/VA. Solid curves: 2 C/min heating rate; double dotted curves: 5 C/min heating rate: dashed curves: 10 C/min heating rate. Perkin
Elmer TGA7. Source: From Arbab, S., Zeinolebadi, A., 2017. Polym. Degrad. Stab. 139, 107. Reprinted with permission from Elsevier. Table 18.3 Mass losses in air related to differential scanning calorimetry (DSC) characteristic temperatures in air and nitrogen for polyacrylonitrile PAN/IA/MA, PAN/IA, and PAN/vinyl acetate (VAc) fibers at various heating rates. Fibers

PAN/IA/MA

DSC atmosphere

Air

N2 PAN/IA

Air

N2 PAN/VAc

Air N2

Characteristic DSC temperatures

Ti Tpk1 Tpk2 Ti Tpk Ti Tpk1 Tpk2 Ti Tpk Ti Tpk Ti Tpk

Mass loss (%) in air at various heating rates 2 C/min

5 C/min

10 C/min

0.05 0.86 3.73 0.05 1.23 0.08 1.2 4.17 0.08 0.97 0.67 2.92 0.67 4.87

0.01 2.7 7.06 0.36 2.98 0.64 3.17 6.85 0.64 2.32 0.64 4.01 0.64 8.3

0.1 3.46 1.34 4.53 0.06 1.90 15.07 0.06 2.13 0.19 29.61 0.19 27.18

Ti, Initiation temperature; Tpk1, first peak temperature; Tpk2, second peak temperature. Source: From Arbab, S., Zeinolebadi, A., 2017. Polym. Degrad. Stab. 139, 107.

308

Thermal Analysis of Textiles and Fibers

temperature at which slope of the mass loses curve drops. This is due to the upward shift of reactions that cause mass losses. At higher heating rates, the free outward diffusion of gases from the sample is inhibited leading to an increase in the decomposition temperature. At lower heating rates, however, the sample temperature is more uniform and diffusion of product gases can occur within the sample, lowering decomposition temperature (Hatakeyama and Quinn, 1999). In general, total mass loss of PAN fibers during stabilization is resultant of the mass loss (due to dehydrogenation and chain rupture) and mass gain (due to addition of oxygen to polymer chains). Some parts of dehydrogenation reactions result from oxidation reactions. In this case, hydrogen leaves the fiber in the form of H2O (Wang, 1998), while other parts of dehydrogenation reactions are independent from oxidation. The higher the heating rate, the more probable the chain rupture and the more intense the mass loss. This is due to simultaneous occurrence of exothermic reactions which makes their control difficult. In addition, diffusion of oxygen into fibers is less than that of higher heating rates due to shorter available times (Huang, 2009). Thus the significant mass loss at heating rate of 10 C/min can be a result of lower oxygen uptake in polymer chains. In nitrogen the mass loss starts earlier, and it is more intense than in air. All this indicated the significance of oxygen absorption in the air measurements (see Fig. 18.7). The shape of the mass loss depends considerably on the chemical composition: for PAN/VAc it is very sudden, but for the other two fibers it is somewhat slower. The heating rate dependence for PAN/IA/MA and PAN/IA is barely noticeable, while for PAN/VAc the types of reactions do change with increasing heating rate. The total mass loss in nitrogen is the sum of dehydrogenation reactions and chain rupture, and mass change due to the oxygen uptake and expel of H2O from structure is not present. From the data of measurements in air at heating rates of 2 and 5 C/min, the total mass losses of PAN/IA/MA, PAN/IA, and PAN/VAc fibers at 400 C are about 13%, 14%, and 17%, respectively (Table 18.4). At a rate of heating of 10 C/min, the total mass loss of PAN/IA/MA, PAN/IA, and PAN/VA fibers shifts to about 30%, 24%, and 43%, respectively (Table 18.4). However, the differences between total mass losses of fibers in nitrogen at heating rates of 2, 5, and 10 C/min are less than 6%. At the same time the difference between the total mass losses of fibers at heating rate of 10 C/min in air and nitrogen is insignificant. Chain rupture may be more intense at a rate of heating of 10 C/min and can be considered as the main factor in drastic mass loss of fibers. According to Table 18.4, the effect of atmosphere on the progress of the reactions causing mass losses is more significant at lower heating rates. Considering limited chain rupture at lower heating rates and better diffusion of oxygen into fibers, the absence of oxygen in measurements in nitrogen may lead to drastic changes in the total mass loss of the fibers. Hence, the mass loss in nitrogen is higher than air due to the absence of mass gain resulting from oxygen uptake (Xiao et al., 2011). The rate of diffusion of oxygen into fibers decreases with an increase in heating rate. The time available for oxygen uptake is also reduced (Hou et al., 2008). Therefore oxygen uptake is limited and the effect on the total mass loss is

Thermal analysis of acrylic and carbon fibers

309

Table 18.4 Total mass losses of polyacrylonitrile (PAN)/IA/MA, PAN/IA, and PAN/vinyl acetate (VAc) fibers at 400 C at various heating rates in air and nitrogen. Fibers

PAN/IA/MA PAN/IA PAN/VAc

Rate 5 2 C/min

Rate 5 5 C/min

Rate 5 10 C/min

Air (%)

N2 (%)

Air (%)

N2 (%)

Air (%)

N2 (%)

13.38 14.47 17.15

24.30 24.99 39.06

13.11 13.26 16.09

22.57 19.99 37.93

29.61 24.06 43.50

27.74 25.26 43.54

Source: From Arbab, S., Zeinolebadi, A., 2017. Polym. Degrad. Stab. 139, 107, Table 3. Reprinted with permission from Elsevier.

Figure 18.5 TMA curves of PAN fibers in air with various heating rates. (A) PAN/IA/MA, (B) PAN/IA, and (C) PAN/VAc. Mettler-Toledo TMA/SDTA840. Source: From Arbab, S., Zeinolebadi, A., 2017. Polym. Degrad. Stab. 139, 107, Fig. 5(a), (b), and (c). Reprinted with permission from Elsevier.

less significant. In addition, more intense chain rupture at higher temperatures leads to drastic mass losses in both air and nitrogen.

TMA measurements TMA curves of PAN/IA/MA, PAN/IA, and PAN/VAc fibers are shown in Fig. 18.5 and the values of shrinkage are given in Table 18.5. Physical, chemical, and total

Table 18.5 Temperature range and extent of physical and chemical shrinkage of polyacrylonitrile (PAN)/IA/MA, PAN/IA, and PAN/vinyl acetate (VAc) fibers and their Tg at various heating rates. PAN/IA/MA

PAN/IA

PAN/VA

Heating rate/ C/min

2

5

10

2

5

10

2

5

10

Physical shrinkage Temp. range ( C) Chemical shrinkage Temp. range ( C) Total shrinkage Tg ( C)

7.0 (70
172) 20.7 (172
274) 27.8 68

7.2 (74
175) 21.2 (175
276) 28.4 69.5

8.9 (76
184) 34.4 (184
280) 43.3 71

10.4 (72
186) 15.2 (186
269) 25.5 60

10.5 (74
188) 15.6 (188
270) 26.1 61.5

11.6 (76
210) 20.2 (210
286) 31.8 72





42.9 72





42.9 78





43 86

Source: From Arbab, S., Zeinolebadi, A., 2017. Polym. Degrad. Stab. 139, 107, Table 4.

Thermal analysis of acrylic and carbon fibers

311

shrinkage of PAN/IA/MA and PAN/IA fibers do not change noticeably with an increase in heating rate from 2 to 5 C/min and the difference between their values is less than 1% (Table 18.5). Increasing the heating rate to 10 C/min shifts the initiation temperature of chemical shrinkage to higher temperatures. In addition, a significant increase in extent of chemical and total shrinkages of PAN/IA/MA and PAN/IA fibers can be observed. This shows the enhancement of cyclization reactions and crosslinking at higher heating rates, leading to more intense chemical shrinkage (Fitzer et al., 1986). However, the increase in physical shrinkage at higher heating rates is insignificant, showing it is not considerably influenced by the heating rate. It is due to the release of physical stress (decrease of orientation), which developed during the spinning and drawing processes (Wang, 1998). The difference between physical shrinkage of PAN/IA/MA and PAN/IA fibers can be attributed to different draw ratios during the fiber preparation process. Greater physical shrinkage of PAN/IA fibers compared to PAN/IA/MA fibers shows the higher stress release during thermal treatment of PAN/IA fibers. This can be related to higher degree of orientation of PAN/IA fibers, which is also proved by sonic modulus. The PAN/IA fibers might have been drawn to higher draw ratios. It is not possible to separate the temperature range of physical and chemical shrinkage in PAN/VAc fibers, because they are overlapped. In addition, there is not a noticeable temperature shift in TMA curves of PAN/VA fibers at different heating rates. Furthermore, the total shrinkage of PAN/VA fibers is unchanged at different heating rates (Table 18.5). Characteristics of PAN/IA/MA and PAN/IA fibers, that is, molecular mass, intrinsic viscosity, sonic modulus, and chemical composition are largely analogous, whereas characteristics of PAN/VA fibers are different from those of PAN/IA/MA and PAN/IA fibers. The difference in orientation of initial PAN fibers might affect thermo-rheological response and thermal behavior of fibers (Seyler, 1994). Thus PAN/VAc fibers have different shrinkage behavior almost independent of heating rate. This shows independency of cyclization reactions and crosslinking from heating rates in this fiber. It is clearly seen that the effect of heating rate on thermal behavior of different fibers is not identical for the three fibers. The glass transition temperature (Tg) of PAN fibers at different heating rates can be measured from TMA curves (Table 18.5). With a change in heating rate, not only chemical reactions but also the physical intermolecular motions are affected, increasing heating rate results in higher Tgs. At higher hating rates, the time available for energy absorption in the chains is reduced.

DMA characterization Okajima et al. (1968) were the first scientists to characterize PAN fibers by DMA on fibers prepared by wet spinning and then drawn in water to 4
16 3 ratio at various temperatures. Finally, they air dried and steam treated these fibers at 80 C
180 C at constant length. They determined the β-transition at 70 C and the αa-transition (amorphous) at 110 C for the dry fibers. The activation energies for these transitions were 84 and 209 kJ/mol, respectively. However, for the steamtreated transitions a new DMA peak was recorded at 160 C. The intensity of this

312

Thermal Analysis of Textiles and Fibers

Figure 18.6 DMA curves (E0 , Ev, and tan δ) for undrawn PAN yarn. Vibron DDV-II viscoelastometer, frequency 110 Hz, heating rate 1 C/min. Source: From Minami, S., 1974. Appl. Polym. Symp. 25, 145, Fig. 1. Reprinted with permission of Wiley.

peak changed parallel to the intensities measured by X-ray, so the authors suggested that this transition is a crystalline relaxation. Minami (1974) found similar peaks by DMA, but they did not see the 70 C transition. However, they observed a new peak at 280 C, and they assigned this peak to cyclization of AN. As described earlier, the nature of the PAN transitions is not clear even today. Minami (1974) attributed the 160 C transition (Figs. 18.6 and 18.7) to the glass transition of PAN, and the 110 C transition to the paracrystalline relaxation. He did this because the intensity of the 160 C peak decreases with increasing chain orientation, and amorphous PAN exhibited one DMA peak only (at 165 C). Ferguson and Ray (1981) found two peaks only for the PAN homopolymer, at 110 C and 280 C. For copolymers of AN with acrylic acid and vinyl pyrrolidone, the intensity of the 110 C peak was decreasing, and a new peak started to grow at 160 C. All these measurements suggest that the 110 C is related somehow to the paracrystalline phase of PAN and the 160 C is the glass transition. Bell and Murayama (1968) obtained a linear relationship between the loss modulus and dye diffusion constant for PAN copolymer fibers (VAc comonomer). Similar results had been found earlier for Nylon 66 fibers (Bell, 1968). Therefore it may be suggested that the interrelation between the dye diffusion constant and the loss modulus is a general phenomenon.

18.5

Oxidative stabilization of polyacrylonitrile fibers

Oxidative stabilization of PAN fibers has two goals: (1) to produce inherently flame-resistant PAN and (2) to make a precursor for CF production. This

Thermal analysis of acrylic and carbon fibers

313

Figure 18.7 Tan delta as a function of temperature for drawn PAN fibers. Curves from top to bottom: undrawn fiber, DR 5 3 3 , DR 5 5 3 , DR 5 7 3 . Vibron DDV-II viscoelastometer, frequency 110 Hz, heating rate 1 C/min. Source: From Minami, S., 1974. Appl. Polym. Symp. 25, 145. Reprinted with permission of Wiley.

stabilization is carried out on the drawn PAN fiber by a heat treatment in air at 220 C
330 C. During this treatment, all the hydrogen bonds in PAN break, the polymer is oxidized and exits the class of thermoplastics.

18.5.1 Effect of chemical composition The chemical composition of the PAN fibers is defined by the type, content, and distribution of comonomers in the chains, and chain characteristics, such as molecular mass and its distribution. More often but not exclusively, PAN homopolymer is used for CF production. PAN-based CFs are produced by the oxidative stabilization of a PAN precursor, normally followed by a two-stage carbonization process. The goal of this procedure is to remove all elements other than carbon. The acrylic precursor is stabilized by controlled heating (220 C
330 C) in air to convert PAN to a form that can be heat treated at higher temperature without the occurrence of melting. The energy release during the oxidative stabilization is considerable: it is about 2000 J/g. Therefore special precautions must be taken to prevent run-away exothermic reactions. In the scientific investigation of this problem (so also in DSC measurements), low heating rates must be used (Morgan, 2005). Industrially, special airflow designs are used to ensure proper dissipation of the developed heat, so as to be able to control the oven temperature. As mentioned already, various reactions take place during the

314

Thermal Analysis of Textiles and Fibers

Figure 18.8 Sematic diagram of the reactions during stabilization process of PAN fibers with different order of occurrence. (A) cyclization of linear PAN, (B) dehydrogenation of linear PAN, (C) dehydrogenation of cyclized PAN, (D) cyclization, and oxidation of PAN, (E) oxidation of cyclized and dehydrogenized PAN. Source: From Arbab, S., Zeinolebadi, A., 2013. Polym. Degrad. Stab. 98, 2537, Fig. 2. Reprinted with permission of Elsevier

oxidative stabilization of PAN, including oxidation, dehydrogenation, and cyclization (Fig. 18.8) (Sen et al., 2003; Rahaman et al., 2007; Arbab and Zeinolebadi, 2013, 2017; Huang, 2009; Gupta et al., 1995). Commercial PAN fibers are commonly prepared with copolymerization of AN and some acidic, amid, and ester comonomers, such as VAc or methyl acrylate, to increase solubility of PAN for the spinning or improve dyeing properties of the fibers. In modacrylics, in addition to the abovementioned comonomers, halogenated vinyl monomers are also present to ensure flame retardancy (Jaffe et al., 1997; He et al., 2005; Morgan, 2005; Shen et al., 2008; Bashir, 1991; Frank et al., 2012; Ouyang et al., 2008a,b, 2009; Wangxi et al., 2003; Rangarajan et al., 2002). The comonomers may also promote the stabilization of PAN fiber, increase the steric hindrance between the chains, and also affect the crystallinity, chain orientation, and the density of the polymer (Crook et al., 2010). Comonomers affect the nature and onset of reactions during the stabilization step, as well as the rate of heat release during stabilization (Morgan, 2005; Frank et al., 2012; Ouyang et al., 2008a,b; Crook et al., 2010; Liu et al., 2009; Zhao et al., 2008). It also affects the steric hindrance of polymer chains and thus alters diffusion rate of oxygen to the chains (Frank et al., 2012; Ouyang et al., 2008a,b; Zhao et al., 2008; Yu et al., 2008; Bajaj et al., 2001). The influence of different comonomers on the initiation temperature of stabilization reactions was studied by Grassie and

Thermal analysis of acrylic and carbon fibers

315

McGuchan (1972). Their results indicated that acidic comonomers facilitate cyclization reactions of the nitrile groups. Itaconic acid is a preferred comonomer due to presence of two carboxylic acid groups that enhance the chance of ionic interaction with nitrile groups of PAN (Ouyang et al., 2008a,b; Bajaj et al., 2001; Grassie and McGuchan, 1972). DSC studies showed that in the presence of acidic comonomers, the exothermic peak is double and broad (Gupta et al., 1995). These AN copolymers are random copolymers, that is, the comonomers are distributed randomly in the chain. It is more difficult to control the sequence composition along the chain when more than one comonomer is used (Ju et al., 2013). The random distribution of the comonomers makes the stabilization reaction more complex. For instance, acidic comonomers change the mechanism of the cyclization reactions of the nitrile groups by initiating it through an ionic mechanism (Ouyang et al., 2008a,b; Rangarajan et al., 2002; Yusof and Ismail, 2012). It is believed that controlling the position of the comonomers and sequence order may improve the efficiency of the cyclization and improve the mechanical properties of CFs (Moskowitz and Wiggins, 2016a). Several researchers used semibatch reactions in conjunction with reversible addition-fragmentation chain transfer polymerization in order to control the distribution of the comonomer units along the chain (Moskowitz and Wiggins, 2016a). The molecular mass of PAN polymers can also influence the morphology, crystallinity, void size, etc., and affect thermal properties of the final CFs (Moskowitz and Wiggins, 2016b). Moskowitz and Wiggins (2016b) claimed that molecular mass and dispersity of polymer chains may have a profound effect on the cyclization behavior of PAN fibers.

18.5.2 Effect of processing parameters Thermal stabilization of PAN fibers is a complicated, time-consuming step in the production of CFs, during which PAN fibers are heated between 230 C and 330 C in oxidizing atmosphere for about 2 hours (Yusof and Ismail, 2012). Processing parameters, such as airflow rate, temperature, tension, heating rate, and heating atmosphere, affect stabilization process of PAN fibers (Zhu et al., 2011; Ouyang et al., 2008a,b; Sabet et al., 2018). These parameters also affect the rate and order of occurrence of the stabilization reactions, as well as the temperature range in which they take place (Rahaman et al., 2007; Arbab and Zeinolebadi, 2013; Huang, 2009; Ju et al., 2014). Various reactions including cyclization, dehydrogenation, and oxidation occur during this step (Sen et al., 2003; Rahaman et al., 2007; Arbab and Zeinolebadi, 2013). The quality of CFs exhibits a significant dependence on the progress of stabilization reactions. These reactions are very complicated and they are influenced by each other. Temperature range, intensity, and order of occurrence of these reactions are affected by heat treatment conditions including atmosphere and heating rate as well as chemical composition of initial PAN fibers (Mirbaha et al., 2013; Rahaman et al., 2007; Arbab and Zeinolebadi, 2013; Huang, 2009; Ju et al., 2014; Arbab et al., 2014; Fitzer et al., 1986; Bajaj et al., 2001). For instance, the heating rate impacts the stabilization time and consequently the production costs

316

Thermal Analysis of Textiles and Fibers

(Fitzer et al., 1986). High heating rates are of interest for the industry from an economic point of view. However, stabilization reactions are exothermic and difficult to control at high heating rates, because it is difficult to dissipate the suddenly developed heat. Kinetics of thermochemical reactions in PAN fibers have been studied by some researchers. It was shown that increasing heating rate enhances chemical reactions and consequently increases heat release during reactions (Ouyang et al., 2008a,b; Wang, 1998). Since stabilization reactions of PAN fibers are highly exothermic, the heating process should be carefully controlled. At higher heating rates, more heat is released in less time. In addition, more damage and defects are introduced to the structure of stabilized fibers. In this condition, fiber burning is likely to occur, limiting the structural development for achieving ideal carbon structure (Fitzer et al., 1986; Hou et al., 2008). Therefore tensile strength and elongation at break of final CFs will be affected (Fitzer et al., 1986). On the other hand, higher heating rates reduce the stabilization time and production costs (Fitzer et al., 1986). Therefore besides the effect of heating rate on the stabilization process, determination of optimum heating rate was also of interest from an economic point of view. According to some researchers, lower heating rates are desirable for stabilization of PAN fibers, leading to better diffusion of oxygen into the core of fibers (Huang, 2009; Zhu et al., 2011). Considering different heating rates from 0.25 to 10 C/min, Hou et al. (2008) observed reduction in density and oxygen content of fibers at higher heating rates. Formation of skin
core structure at higher heating rates was attributed to lower diffusion of oxygen into core of fibers as a result of shorter reactions times (Hou et al., 2008). Furthermore, controlling the structural changes toward formation of intermediate cyclic structure becomes more difficult at higher heating rates due to the increase in rate of interactions. There is mass loss in PAN fibers, due to expel of volatile by-products during thermal stabilization. The total mass loss of PAN fibers is the sum of mass loss and mass gain. Mass loss results from elimination of hydrogen, HCN, NH3, H2O, etc. Chain rupture is another reason for mass loss of fibers during stabilization process. Polymer chains in the fiber surface layer are more readily exposed to heat. Thus they are more vulnerable to damage and rupture, compared to the chains in the core of the fibers. Mass gain during the heating process is due to oxidation reactions (formation of carbonyl and hydroxyl groups). However, hydroxyl groups may also participate in the dehydrogenation reactions at higher temperatures and contribute to the total mass loss (Sen et al., 2003; Bajaj et al., 2001). It is not easy to separate various interactions contributing to the mass loss. No accurate mechanism has been proposed for their separation. The heating rate determines the release rate of volatile components from PAN fibers and consequently affects the structure formation and properties of final fibers (Yusof and Ismail, 2012; Zhu et al., 2011). Since PAN fibers are bad conductors of heat, acceleration of exothermic reactions at higher heating rates might lead to overheating in some parts of the chains and consequently chain rupture and rapid mass loss would take place (Hanna et al., 2012; Liu, 2010). It has been verified that applying dimensional constraint or stretch (tension) is an effective way to prevent the macromolecular chains from entropic relaxation or

Thermal analysis of acrylic and carbon fibers

317

physical shrinkage. Accordingly, thermal stretching prior to oxidative stabilization develops the orientation degree of PAN chains, whereas applying stretch or dimensional constraint during chemical shrinkage may hinder cyclization reactions during stabilization of PAN fibers and make it difficult for nitrile groups to form ladderlike structures (Sabet et al., 2016; Lian et al., 2012). As a result, the structural imperfections are likely to be retained in the final CFs, deteriorating their mechanical properties (Lian et al., 2012). Furthermore, studies of Hou et al. (2008) revealed that increasing the heating rate up to a specific rate decreases total shrinkage of PAN fibers, while at even higher rates, the total shrinkage increases with increasing heating rate. According to their research, the extent of cyclization reactions was decreased with an increase in heating rate, whereas the extent of intermolecular crosslinking was increased (Asakawa et al., 2012). Fitzer et al. (1986) showed that with an increase in heating rate, the initiation of chemical shrinkage shifts to higher temperatures. Further increase in the heating rate above a specific value shifts the initiation of chemical shrinkage to lower temperatures. During stabilization step of PAN fibers, various reactions including cyclization, dehydrogenation, and oxidation can take place in the fibers causing an exothermic peak in DSC curves (Fig. 18.2) (Sen et al., 2003; Rahaman et al., 2007; Arbab and Zeinolebadi, 2013; Gupta et al., 1995). The stabilization reactions of PAN fibers in air and nitrogen are different. Cyclization, dehydrogenation, and oxidation reactions can take place in air, but oxidation reactions are mostly absent in nitrogen (Sen et al., 2003). Dehydrogenation reactions can be divided into two categories. The first category is dehydrogenation reactions resulting from oxidation that causes hydrogen to leave the fiber in the form of H2O. The second category is dehydrogenation reactions that are independent from oxidation (Arbab and Zeinolebadi, 2013). Dehydrogenation reactions resulting from oxidation do not take place in nitrogen (Rahaman et al., 2007; Arbab and Zeinolebadi, 2013). Hence, it is possible to separate oxidation from other reactions by comparing DSC curves recorded in air and nitrogen (Arbab and Zeinolebadi, 2013, 2017). Mass loss of PAN fibers during the heating process is a result of dehydrogenation reactions and chain rupture, whereas cyclization reactions do not lead to significant mass loss. Comparing the results of TGA and DSC measurements, one can determine the temperature range of cyclization, oxidation, and dehydrogenation reactions (Arbab and Zeinolebadi, 2013, 2017). Moreover, determination of the temperature range of chemical shrinkage can help to confirm the temperature range of cyclization.

18.5.3 Study of the stabilization by Fourier-transform infrared spectroscopy (FT-IR) It is not certain how the chemical composition of the PAN fiber can affect the stabilization reaction, because too many variations can be found in commercial PAN fibers. In this section, we will try to focus on the effect of chemical composition on the thermal response during the stabilization step. As an example, we studied how changes in the chemical composition of different type of PAN fibers may affect the stabilization.

318

Thermal Analysis of Textiles and Fibers

Table 18.6 Assignment of IR bands for polyacrylonitrile fibers. Wavenumber (cm21)

Functional groups and mode of vibration

Wavenumber (cm21)

Functional groups and mode of vibration

2243 1595

C  N (stretching) Combination of C 5 N, C 5 C (stretching), N
H (in-plane bending) C 5 O (stretching) CH2 (in-plane bending) CH2 (stretching)

1224 1250

CH2 (twisting) CH2 (bending)

1073 3450 1100
1300

CH (bending) OH (stretching) C
O (stretching)

1730 1450 2932

Figure 18.9 FT-IR spectra of initial and stabilized PAN/IA/MA fibers at different stabilization temperatures. Source: From Arbab, S., Zeinolebadi, A., 2013. Polym. Degrad. Stab. 98, 2537. FTIR Nicolet Magna IR560F. Reprinted with permission from Elsevier.

Table 18.6 summarizes the assignment of IR bands recorded for PAN fibers. Fig. 18.9 shows changes in FT-IR spectra for PAN/IA/MA (AN/itaconic acid/ methyl acrylate copolymer) fibers. As the stabilization temperature increases, the intensity of the C  N stretching band at 2243 cm21 decreases above 200 C. At the same time, a new band appears at 1595 cm21 which can be attributed to the combination effect of C 5 C, C 5 N, and N
H groups. With further increase of the stabilization temperature to 225 C, the intensity of this band grows significantly evincing the increase of the number of these groups in the structure of stabilized fibers. These changes confirm the gradual formation of cyclized structure in PAN fibers during the stabilization. At 250 C, the C  N band is still present in the FTIR spectra. This indicates that the stabilized fiber at 250 C is still not fully cyclized.

Thermal analysis of acrylic and carbon fibers

319

Figure 18.10 FT-IR spectra of initial and stabilized PAN/IA fibers at different stabilization temperatures. FT-IR Nicolet Magna IR560F. From Arbab and Zeinolebadi (2013). Source: From Arbab, S., Zeinolebadi, A., 2013. Polym. Degrad. Stab. 98, 2537, Figs. 3 and 4. Reprinted with permission of Elsevier.

At temperatures higher than 200 C, the intensity of the 1730 cm21 band, attributed to the C 5 O stretching in acidic (IA) and ester (MA) comonomers, is decreasing. At 250 C this peak completely disappears. At the same time a shoulder-like peak attributed to the presence of the C 5 O groups in the cyclized PAN shows up at 1715 cm21 (Ouyang et al., 2008a,b; Wangxi et al., 2003; Lee, 2012). According to theoretical predictions, the C 5 O groups from a comonomer initiate the cyclization reactions and crosslink the structure with ionic mechanism (Wangxi et al., 2003; Rangarajan et al., 2002; Sivy and Coleman, 1981). Therefore with the progress of cyclization reactions, the number of the C 5 O groups in PAN decreases and increases in the cyclized structure due to the oxidation reactions. The intensity of stretching and bending vibrations of aliphatic methylene at 2932 and 1450 cm21 also decreases. This can be attributed to the conversion of linear structure to cyclized structure and reduction of the number of aliphatic methylene groups at higher stabilization temperatures. In summary, FT-IR spectra at different stabilization temperatures show that the major changes in the chemical structure of PAN/ IA/MA fibers occur above 225 C. FT-IR spectra of PAN/IA (AN/itaconic acid copolymer) fibers show similar changes during the stabilization process, see Fig. 18.10. These changes include a decrease in the intensity of the C  N stretching band at 2243 cm21 above 225 C, and simultaneous growth of a band at 1595 cm21 attributed to simultaneous effect of the C 5 C, C 5 N, and N
H groups. Other changes in the FT-IR spectra of PAN/IA fibers include a decrease in intensity of the 2926 and 1450 cm21 bands assigned to stretching and bending vibrations of aliphatic methylene groups. These changes confirm the conversion of linear structure of PAN to a partially cyclized

320

Thermal Analysis of Textiles and Fibers

Figure 18.11 FT-IR spectra of raw and stabilized PAN/VAc fibers at different stabilization temperatures. FT-IR Nicolet Magna IR560F. Source: From Arbab, S., Zeinolebadi, A., 2013. Polym. Degrad. Stab. 98, 2537, Fig. 4. Reprinted with permission from Elsevier.

structure containing C 5 O, C 5 C, C 5 N, and N
H groups in cyclized parts and C  N groups in linear parts of the chains (Wangxi et al., 2003; Rangarajan et al., 2002; Kakida and Tashiro, 1997; Duan et al., 2012). The presence of itaconic acid comonomers in PAN/IA fibers initiates reactions with ionic mechanism. As a result, a decrease of the intensity of the 1730 cm21 band (C 5 O groups in acidic comonomer) and appearance of a shoulder-like peak at 1715 cm21 (C 5 O groups in cyclized structure) can be observed. FT-IR spectra of PAN/VAc fibers in Fig. 18.11 show that at stabilization temperature of 250 C, the intensity of the band at 2243 cm21 (C  N groups) and the band at 2932 cm21 (aliphatic methylene groups) did not decrease significantly. The band at 1732 cm21 (C 5 O groups in VAc comonomers) did not change and the band at 1595 cm21 attributed to the formation of cyclized structure did not grow significantly compared to other fibers. In general, the changes related to formation of cyclized structure in PAN/VAc at 250 C are less than in the PAN/IA/MA and PAN/IA fibers. This reveals lower conversion of the linear structure to partially cyclized structure in PAN/VAc than in other fibers. According to a review published by Bashir (1991), there are four major interpretations of dominant reactions occurring during stabilization of PAN fibers: (1) intramolecular nitrile cyclization with the formation of cyclized C 5 N containing structures (ladder polymer), (2) intermolecular nitrile reaction giving crosslinking, (3) azomethine crosslinking, and (4) hydrogen elimination leading to the formation of conjugated C 5 C bonds in the backbone of chains. These reactions increase compactness and density of PAN (Wangxi et al., 2003). Despite numerous studies (Kakida and Tashiro, 1997; Sen et al., 2003; Rahaman et al., 2007; Grassie and McGuchan, 1972; Shimada and Takahagi, 1986), there is no agreement on the order of occurrence of stabilization reactions. On the basis of

Thermal analysis of acrylic and carbon fibers

321

different analytical methods, researchers presented different interpretations regarding the order of occurrence of these reactions. FT-IR and elemental analysis studies by Shimada and Takahagi (1986) showed that the three stabilization reactions (i.e., cyclization, dehydrogenation, and oxidation) occur almost simultaneously. By using a simultaneous DSC/FT-IR method, Gupta et al. (1995) attributed the first exothermic peak to cyclization and dehydrogenation and the second peak to oxidation reactions. By isothermal DSC/FT-IR spectroscopy on PAN homopolymer and AN /methacrylic acid copolymer (AN/MAA), Kakida and Tashiro (1997) found that dehydrogenation is the dominant reaction, and the MAA comonomer has a greater impact on dehydrogenation reactions, compared to cyclization. Sen et al. (2003) suggested that thermal reactions occur in the following order: first cyclization, then dehydrogenation, and finally oxidation. This discrepancy originates mainly from the differences in the chemical composition of the studied fibers, which is one of the major factors determining the mechanism of reactions (Rangarajan et al., 2002; Bajaj et al., 2001). It is worth mentioning here that a correct map of stabilization reactions is essential for adjusting the processing conditions (such as temperature and tension) in order to obtain special CFs with high mechanical properties such as T700 and T1000.

References Andrews, R.D., Kimmel, R.M., 1965. Polym. Lett 3, 167. Arbab, S., Nourpanah, P., Mohammadi, N., Soleimani, M., 2008. J. Appl. Polym. Sci. 109, 3461. Arbab, S., Nourpanah, P., Mohammadi, N., Zeinolebadi, A., 2011a. Polym. Bull. 66, 1267. Arbab, S., Nourpanah, P., Mohammadi, N., Zeinolebadi, A., 2011b. J. Polym. Res. 18, 1343. Arbab, S., Mirbaha, H., Zeinolebadi, A., Nourpanah, P., 2014. J. Appl. Polym. Sci. 131, 403. Arbab, S., Zeinolebadi, A., 2013. Polym. Degrad. Stab. 98, 2537. Arbab, S., Zeinolebadi, A., 2017. Polym. Degrad. Stab. 139, 107. Asakawa, H., Nishida, K., Yamamoto, J., Inoue, R., Kanaya, T., 2012. Polymer 53, 2777. Assignment of the Glass Transition. In: Seyler, R.J. (Ed.), ASTM Publication Code Number (PCN) 04-012490-50. ASTM, Philadelphia. Bacon, R., Filamentary Graphite and Method for Producing the Same, US Patent 2957756, 1960. Bajaj, P., Sreekumar, P.T.V., Sen, K., 2001. Polym. Bull. 42 (4), 1707. Bashir, Z., 1991. Carbon 29 (8), 1081. Bell, J.P., 1968. J. Appl. Polym. Sci. 12, 627. Bell, J.P., Murayama, T., 1968. J. Appl. Polym. Sci. 12, 1795. Berger, W., Heller, A., Thamke, W., Adler, H.-J., 1973. Faserforsh. Textiltech 24, 484. Bohn, C.R., Schaefgen, J.R., Statton, W.O., 1961. J. Polym. Sci. 55, 531. Catta Preta, I.F., Sakata, S.K., Garcia, G., Zimmermann, J.P., Galembeck, F., Giovedi, C., 2007. J. Therm. Anal. Calorim. 87 (3), 657. Chen, J., Wang, C.G., Dong, X.G., Liu, H.Z., 2006. J. Polym. Res. 13, 515. Crook, V., Ebdon, J., Hunt, B., Joseph, P., Wyman, P., 2010. Polym. Degrad. Stab. 95 (12), 2260. Duan, Q., Wang, B., Wang, H., 2012. J. Macromol. Sci., Phys. 51 (12), 2428. Dunn, P., Ennis, B.C., 1970. J. Appl. Polym. Sci. 14, 1795.

322

Thermal Analysis of Textiles and Fibers

Fitzer, E., Frohs, W., Heine, M., 1986. Carbon 24, 387. Fox, T.G., Lashaek, S., 1955. J. Polym. Sci. 15 (80), 371. Frank, E., Hermanutz, F., Buchmeiser, M., 2012. Macromol. Mater. Eng. 297, 493. Frushour, B.G., Knorr, R.S., 1985. In: Lwein, M., Pearce, E.M. (Eds.), Handbook of Fiber Science and Technology, vol. 4. Dekker, New York and Basel, p. 171. Grassie, N., McGuchan, R., 1972. Eur. Polym. J. 8 (2), 257. Gupta, A.K., Paliwal, D.K., Bajaj, P., 1995. J. Appl. Polym. Sci. 58 (7), 1161. Hanna, S.B., Yehia, A.A., Ismail, M.N., Khalaf, A.I., 2012. J. Appl. Polym. Sci. 123, 2074. Hatakeyama, T., Quinn, F.X., 1999. Thermal Analysis
Fundamentals and Applications to Polymer Science. John Wiley and Sons Ltd, Chichester. He, D., Wang, C., Bai, Y., Zhu, B., 2005. J. Mater. Sci. Technol. 21 (3), 376. Hou, Y., Sun, T., Wang, H., Wu, D., 2008. Text. Res. J. 78, 806. Howard, W.H., 1961. J. Appl. Polym. Sci. 5 (15), 303. Huang, X., 2009. Fabrication and properties of carbon fibers. Materials 2, 2369
2403. Jaffe, M., Menczel, J.D., Bessey, W.E., 1997. Chapter 7, Fibers. In: Turi, E.A. (Ed.), Thermal Characterization of Polymeric Materials, second ed. Academic Press. Ju, A., Guang, S., Xu, H., 2013. Carbon 54, 3232018. Ju, A., Luo, M., Zhang, H., Ge, M., 2014. J. Therm. Anal. Calorim. 117, 205. Kakida, H., Tashiro, K., 1997. Polym. J. 29 (4), 353. Krigbaum, W.R., Tokita, N., 1960. J. Polym. Sci. 43, 467. Lee, S., 2012. Adv. Chem. Eng. Sci. 2 (2), 275. Lian, F., Liu, J., Ma, Z., Liang, J., 2012. Carbon 50, 488. Liu, J., Zhou, P., Zhang, L., Ma, Z., Liang, J., Fong, H., 2009. Carbon 47, 1087. Liu, Y., 2010. Stabilization and Carbonization Studies of Polyacrylonitrile/Carbon Nanotube Composite Fibers. Georgia Institute of Technology. Masson, J., 1995. Acrylic Fiber Technology and Applications, Marcel Dekker, New York. Menczel, J., Wunderlich, B., 1981. J. Polym. Sci., Polym. Lett. Ed. 19, 261. Menczel, J.D., 2020, to be published. Minami, S., 1974. Appl. Polym. Symp. 25, 145. Mirbaha, H., Arbab, S., Ahmad, Z., Nourpanah, P., 2013. Smart Mater. Struct. 22 (4), 045019. Mittal, J., Mathur, R.B., Bahl, R.B., Inagaki, M., 1998. Carbon 36, 893. Morgan, P., 2005. Carbon Fibers and Their Composites. Taylor & Francis Group, CRC Press, Boca Raton, FL. Moskowitz, J.D., Wiggins, J.S., 2016a. Polymer 84, 311. Moskowitz, J.D., Wiggins, J.S., 2016b. Polym. Degrad. Stab. 125, 76. Okajima, S., Ikeda, M., Takeuchi, A., 1968. J. Polym. Sci., Part A- 1 (6), 1925. Ouyang, Q., Cheng, L., Wang, H., Li, K., 2008a. Polym. Degrad. Stab. 93 (8), 1415. Ouyang, Q., Cheng, L., Wang, H., Li, K., 2008b. J. Therm. Anal. Calorim. 94 (1), 85. Ouyang, Q., Cheng, L., Wang, H., Li, K., 2009. e-Polymers 9 (1), 015. Paiva, M.C., Kotasthane, P., Edie, D.D., Ogale, A.A., 2003. Carbon 41, 1399. Rahaman, M.S.A., Ismail, A.F., Mustafa, A., 2007. Polym. Degrad. Stab. 92, 1421. Rangarajan, P., Yang, J., Bhanu, V., Godshall, D., McGrath, J., Wilkes, G., et al., 2002. J. Appl. Polym. Sci. 85, 69. Sabet, E.N., Nourpanah, P., Arbab, S., 2016. Polymer 90, 138. Sabet, E.N., Nourpanah, P.S., Arbab, S., 2018. Adv. Polym. Technol. in press. Salem, D., 2001. Structure Formation in Polymeric Fibers. Hanser, Munich. Sa´nchez-Soto, P.J., Aviles, M.A., del Rio, J.C., Gines, J.M., Pascual, J., Pe´rez-Rodriguez, J. L., 2001. J Anal. Appl. Pyrolysis 58-59, 155.

Thermal analysis of acrylic and carbon fibers

323

Schick, C., Androsch, R., Schmelzer, J.W.P., 2017. J. Phys.: Condens. Matter 29, 453002 (35pp). Sen, K., Bajaj, P., Sreekumar, T., 2003. J. Polym. Sci., B: Polym. Phys. 41, 2949. Shen, X., Ji, Y., Wang, J., 2008. J. Appl. Polym. Sci. 110 (1), 313. Shimada, I., Takahagi, T., 1986. J. Polym. Sci., A: Polym. Chem. 24, 1989. Sivy, G.T., Coleman, M.M., 1981. Carbon 19 (2), 127. Slade, P.S., 1970. Thermochim. Acta 1, 5459. Sviridov, A.A., Varshavskii, V.Y., Seleznev, A.N., Morozov, V.A., Kepman, A.V., 2009. Fibre Chem. 41 (4), 236. 2009. Thu¨nemann, A.F., Ruland, W., 2000. Macromolecules 33, 1848. Wang, P.H., 1998. J. Appl. Polym. Sci. 67, 1185. Wang, Y.X., Wang, C.G., Yu, M., 2007. J. Appl. Polym. Sci. 104, 3723. Wangxi, Z., Jie, L., Gang, W., 2003. Carbon 41 (14), 2805. Warner, S.B., Uhlmann, D.R., Peebles Jr, L.H., 1979. J. Mater. Sci. 14 (8), 1893. Xiao, S., Lv, H., Tong, Y., Xu, L., Chen, B., 2011. J. Appl. Polym. Sci. 122, 480. Yu, M.J., Wang, C.G., Bai, Y.J., Zhu, B., Ji, M., Xu, Y., 2008. J. Polym. Sci., B: Polym. Phys. 46, 759. Yusof, N., Ismail, A., 2012. J. Anal. Appl. Pyrolysis 93, 1. Zhao, Y.Q., Wang, C.G., Bai, Y.J., Chen, G.W., Jing, M., Zhu, B., 2008. J. Colloid Interface Sci. 329 (1), 48. Zhu, J., Wei, S., Rutman, D., Haldolaarachchige, N., Young, D.P., Guo, Z., 2011. Polymer 52, 2947.

Further reading Chen, J.C., Harrison, I., 2002. Carbon 40, 25. Hinrichsen, G., 1971. Angew. Makromol. Chem. 20, 121. Kharatyan, S.L., Chatilyan, H.A., Mukasyan, A.S., Simonetti, D.A., Varma, A., 2005. Am. Inst. Chem. Eng. 51, 261. Mathur, R.B., Bahl, O.P., Matta, V.K., Nagpal, K.C., 1988. Carbon 26, 295. Patrick, D.R., Fardo, S.W., 2009. Industrial Process Control Systems. The Fairmont Press, Inc. Sabet, E.N., Nourpanah, P., Arbab, S., 2017. Adv. Polym. Technol. 36, 424. Wunderlich, B., 2005. Thermal Analysis of Polymeric Materials. Springer, Berlin.

Thermal analysis of liquid crystalline polymers

19

Michael Jaffe1, Anthony J. East2 and Xianhong Feng3 1 New Jersey Innovation Institute, Newark, NJ, United States, 2Consultant, Madison, NJ, United States, 3Becton Dickinson and Company, Franklin Lakes, NJ, United States

Abstract Early in the development of polymer science, Prof. H. Mark suggested that the tensile modulus of polymers should correlate with both the chemical and physical structures of the macromolecule. It was further recognized that maximum property levels would be achieved when all of the molecular chain backbone bonds of the polymer were lined up in the direction of measurement. Such an extended chain morphology has been demonstrated with gel spun polyethylene and with nematogenic polyamides and polyesters. Only mainchain liquid crystalline polymers (LCPs) exhibiting nematic behavior in the fluid state have found fiber applications. All of the LCPs are composed of stiff, highly aromatic monomers and are characterized by domains of high local orientation in the solid state (orientation function .0.95). If processed into fibers, the locally oriented domains are transformed into a single domain of high global molecular orientation parallel to the fiber direction. The thermal analysis literature, associated with lyotropic and thermotropic LCP fibers, is not extensive and work through the first decade of the 21st century and is well summarized in the books edited by Turi and by Menczel and Prime. In this chapter, the materials science of LCP fibers is reviewed, the recent application of TA techniques to LCP fibers is summarized, and the utility and impact of thermal analysis techniques in the understanding of LCP process
structure
property relations are discussed in detail.

19.1

Introduction

Early in the development of polymer science, Prof. H. Mark (Guth and Mark, 1934; Mark, 1936) suggested that the tensile modulus of polymers should correlate with both the chemical and physical structures of the macromolecule. It was further recognized that maximum property levels would be achieved when all of the molecular chain backbone bonds of the polymer were lined up in the direction of measurement. Such an extended chain morphology has been demonstrated with gel spun polyethylene (Smith and Lemstra, 1980; Kavesh et al., 2006) and with nematogenic polyamides and polyesters (Jaffe et al., 2018). It is known that polymers can exhibit all the mesophases (liquid crystalline phases) associated with low molar mass molecule liquid crystals (Jaffe, 1987). Polymer liquid crystalline behavior can be a function of main-chain or side-chain chemistry, and polymers may exhibit liquid crystalline properties in the melt (thermotropic) or in solution (lyotropic) or in both. Thermal Analysis of Textiles and Fibers. DOI: https://doi.org/10.1016/B978-0-08-100572-9.00019-7 © 2020 Elsevier Ltd. All rights reserved.

326

Thermal Analysis of Textiles and Fibers

For detailed discussions of the chemistry, physics, and applications of liquid crystalline polymers (LCPs), see the extensive open (Jaffe, 1987, 1991, 2018) and the extensive patent literature (Dupont, Celanese). Only main-chain LCPs exhibiting nematic behavior in the fluid state have found fiber applications. All of the LCPs are composed of stiff, highly aromatic monomers and are characterized by domains of high local orientation in the solid state (orientation function .0.95). If processed into fibers, the locally oriented domains are transformed into a single domain of high global molecular orientation parallel to the fiber direction. LCP fibers are characterized by very high-specific tensile properties (property divided by density) when compared to metals or ceramics. The highly anisotropic nature of oriented LCP fibers causes inherent weakness in shear and compression, limiting their use almost exclusively to applications in tension. It should also be noted that in the absence of global orientation, the tensile properties of the thermotropic polymers are similar to those of filled plastics. LCP fibers were extensively investigated in the 1960s and 1970s, with the lyotropic poly(p-phenyleneterephthalamide) (Kevlar, p-aramid) developed by DuPont (Kwolek, 1966; Hoegger, 1971) and a series of thermotropic all-aromatic polyesters explored by DuPont, Celanese, and others (Calundann and Jaffe, 1982; Jackson, 1982). The all-aromatic main-chain lyotropic polyamides and thermotropic polyesters are semirod-like molecules that are naturally organized into nematic liquid crystal domains when in the fluid state. The nematic state can be viewed as similar to “logs floating in a river,” leading to ease of flow parallel to the molecular axis (low elongational viscosity at high molecular weight) and an extended chain structure in the solid state. Hence, the nematic state in polymers brings both processing ease and high axial tensile properties. Fig. 19.1 illustrates the processing of LCPs and the morphology produced, in contrast to conventional polymers, such as Poly(ethyleneterephthalate) (PET) or nylon.

19.2

Chemical structure of liquid crystalline polymers

Nematic polymer backbone chemistry is characterized by a high degree of aromaticity, planarity, and linearity along the molecular chain. Most common moieties are p-phenylene, 1,4-biphenyl, and 2,6-naphthalyl linked by ester or amide linkages. The all-aromatic polyester homopolymers tend to be intractable, decomposing at temperatures well below their melting points, and are insoluble in almost all solvents. Successful melting point
reduction strategies include incorporation of comonomers to reduce crystal packing, decreasing chain linearity, and/or increasing interchain distances with bulky side groups. Fig. 19.2 illustrates the melting point (X!N transition) of the phenyl/naphthyl thermotropic copolymers as a function of naphthyl content. All of these approaches lower the polymer transition temperature to the melt (or increase solubility) and when the melting temperature is reduced to below the polymer decomposition temperature, stable melt processing is possible. These approaches have led to large numbers of melt-processable thermotropic polyesters,

Thermal analysis of liquid crystalline polymers

327

Conventional (PET)

Liquid crystal

Solution or melt

Random coil

Nematic structure Extrusion

Solid state

Extended chain structure

Lamellar structure

• High chain continuity • High mechanical properties

• Low chain continuity • Low mechanical properties

Figure 19.1 LCP versus PET (Jaffe, 1991, Figure 1). HNA O OC

HQ

TA

HBA

OC

OC

O

CO

O

O

O

OC

I

II Melting point (ºC) 340

II

320 300 280 260 240

I 20 40 60 80 HNA (mol.%)

Figure 19.2 How it works (Jaffe, 1991, Figure 19.3).

well defined in an extensive patent literature (see, e.g., Calundann et al., 1976) and summarized in a number of review papers (Jaffe, 1987; Jaffe et al., 2018). The chemical structure of the most important LCP fibers is shown in Fig. 19.3. Much of

328

Thermal Analysis of Textiles and Fibers

Figure 19.3 Chemistry of p-aramid (Jaffe and Jones, 1985) and Vectran (Jaffe, 1991).

the cost of LCP fibers is the result of high monomer cost and limited monomer availability. There have been no new LCP fibers introduced since the 1980s, although the work of Sekema with “M5” (Lammers et al., 1998) and Dingmans (Knijnenberg et al., 2006) offers promise for improved LCP fibers in the future.

19.3

Process
structure
property relationships of liquid crystalline polymer fibers

All of the nematic lyotropic LCPs are processed by dry jet wet spinning. The process differs from conventional wet spinning in that the spinneret is mounted above the coagulation bath, allowing extrusion to take place at a higher temperature than coagulation and permitting significant fiber drawdown to occur in the air gap between spinneret and bath (Jaffe and Jones, 1985). While lyotropic aromatic-heterocyclic polymer fibers have been explored (Calundann et al., 1988), most recently M5 by DuPont (Lammers et al., 1998), only p-aramid fibers (Kevlar, Twaron) have achieved major commercial success. p-Aramid fibers are dry jet wet spun from a nematic lyotropic solution (100% sulfuric acid solvent), drawn to high molecular orientation in the air gap and taken up after extensive washing. Thermotropic polyesters are melt spun from the nematic phase and orient easily in an elongational flow field to form a highly oriented, extended chain morphology in the solid state. The thermotropic polyester fibers show an initial modulus close to theory—typical values range from 70 to 150 GPa, similar to the values achieved with the lyotropic p-aramids. Ward (Richardson and Ward, 1970) and Northolt (Baltussen and Northolt, 1996) have shown that the tensile modulus of LCP fibers may be described by an “aggregate model,” that is, the modulus is a function of the inherent chain modulus, the molecular chain orientation, and the shear modulus (which describes the stress transfer between chains). Tensile strength of LCP fibers follows the prediction of the “lagshear model” as shown by Yoon (Jaffe, 1987). Both the aggregate and the lag-shear models treat the LCP as though it was a self-reinforced short fiber-reinforced composite. As-spun tensile strength of the thermotropic copolyesters tends to be on the order of about 1 GPa and can be advanced to about 3 GPa by annealing free to shrink close to the melting temperature. Kinetics of strength improvement follows those of solid-state polymerization, leading many researchers to associate strength increases

Thermal analysis of liquid crystalline polymers

329

Table 19.1 Typical fiber mechanical properties. Fiber

Tensile modulus (GPa)

Shear modulus (GPa)

Tensile strength (GPa)

Composite compression strength (GPa)

p-Aramid Thermotropic copolyester PBO M5 Carbon-HM370

70
130 70
130

1.8 1.3

3.2 3.2

240
290 100
200

240 285 370

5.2 17.5

3.5 3.5 2.2

500 700
900

with molecular weight increase (Jaffe, 1987). An alternative, or perhaps complementary mechanism, of strength/toughness enhancement, treating the LCP as an entangled yarn of LCP molecules, has been advanced by Collins (Collins and Long, 1994). Strength and elongation to failure increase in tandem during thermotropic LCP fiber heat treatment, suggesting a mechanism of flaw reduction. Structural perfection and improved intermolecular bonding also play a role in the observed property improvement. Modulus increase during annealing is usually minimal with the thermotropic polyesters, unless structural perfection leads to an increase in overall molecular chain orientation. Lyotropic fibers show high strength directly after spinning and the modulus can be increased by heat treating under tension to improve molecular orientation and crystalline perfection. The unifying feature of all fibers spun from LCPs is very high axial molecular orientation, which leads to extreme anisotropy of microstructure and mechanical properties. In the transverse direction the strength is only about 20% of the axial strength and the modulus is typically less than 10% of the axial value. The microstructure of LCP fibers reflects their very high orientational molecular anisotropy and may be described as a hierarchy composed of fibrillar structures ranging in diameter from microns to about 10 nm (Sawyer and Jaffe, 1986). The properties of the most important LCP fibers are listed in Table 19.1. Key application area for LCP fibers include hard armor (vehicles, helmets), soft ballistic protection (vests), cut protection (gloves), and a variety of composite uses that include honeycomb structure, pressure vessels, and rubber (tire) reinforcement. LCP ropes and cables find utility in the mooring of huge offshore structures, such as oil drilling platforms and the reinforcement and support of optical cables. LCP fibers also find specialty niche markets, such as sails for racing yachts and specialized fishing nets.

19.4

Thermal analysis of liquid crystalline polymer fibers

The thermal analysis literature, associated with lyotropic LCP fibers, p-aramid fibers, and aromatic-heterocyclic fibers, is not extensive and work through the first

330

Thermal Analysis of Textiles and Fibers

Figure 19.4 Thermogravimetric response (10K/min) of Kevlar—Q600 TGA, heating rate 10K/min (Comet et al., 2009, Figure 1).

decade of the 21st century and is well summarized in the books edited by Turi (Jaffe, 1981, 1997) and by Menczel and Prime (2009). All of the lyotropic LCP fibers decompose before melting under ordinary conditions, but multiple thermal analytic technique studies by Brown and Ennis (1977), Penn and Larsen (1979), and Jaffe and Jones (1985) on poly(p-aramid) fibers suggest an effective Tg of B360 C and melting above 500 C, coupled with chemical decomposition. There are several studies of aramid fiber thermal stability where thermo Gravimetric analysis (TGA), differential scanning calorimetry (DSC), and various supporting analytical techniques are utilized. Yang et al. (2011) carried out an extensive TGA study of Kevlar pyrolysis using fourier transform infra-red spectroscopy (FTIR) to analyze products formed. They show that initial Kevlar weight loss is the elimination of bound water, followed by chain scission and the decomposition of the produced p-aramid chain fragments to low molar mass gases, water, CO2, CO, and NH3. Comet et al. (2009) also used a combination of TGA and FTIR to determine the best conditions to sublimate Kevlar fibers to produce Kevlar nanoparticles. They determined the best conditions for Kevlar sublimation between 400 C and 450 C under inert conditions. This conclusion was based on their TGA studies of Kevlar under inert and oxidative conditions, shown in Fig. 19.4. Cai and Yu (2009) studied the thermal degradation kinetics of Kevlar 29, Kevlar 49, poly(benzoxazole) (PBO), and Nomex. Kevlar 29 and 49 and p-aramid variants differ only in tensile modulus. PBO is a lyotropic aromatic-heterocyclic fiber and Nomex is a conventional polymer (m-aramid) wet spun into fiber for flammability protection. Fig. 19.5 shows the TGA and differential TGA response of these fibers. They conclude that Nomex is the most stable at the high temperature, although its degradation starts at the lowest temperature—see comparison of Nomex and Kevlar in Fig. 19.5. In a similar study, Liu and Yu (2006) and Perepelkin et al. (2005) utilized TGA to determine the decomposition behavior of high-performance fibers, including the p-aramids and PBO. In an extensive review of the thermal

Thermal analysis of liquid crystalline polymers

331

Retention of storage modulous (%)

Figure 19.5 TGA of high performance fibers, TG 209 F1 Iris device, 20 C/min heating rate (Cai and Yu, 2009, Figure 1). As received 60/12 week 70/12 week 90/2 week 90/4 week

100

Tan θ

90 As received 60/12 week 70/12 week 90/2 week 90/4 week

80

70 50

100

150 200 250 Temperature (ºC)

300

350

50

100

150 200 250 300 Temperature (ºC)

350

Figure 19.6 Changes in the storage modulus and tan δ, TA Instruments TA100 DMA, 1.0 Hz, 5 C/min (Li et al., 2013, Figures 6 and 7).

stability of high-performance fibers, including all available aromatic-heterocyclic compositions, Bourbigot and Flambard (2002) utilize TGA and flammability testing to conclude that PBO fibers offer the most protection from fire. Similar TGA studies of p-aramid fibers were conducted by Perepelkin et al. (2005) and Druzhinina et al. (2012). Li et al. (2013) utilized dynamic mechanical analysis (DMA) and Small Angle X-Ray Scattering to study the hydrothermal aging of aramid fibers, specifically assessing the changes in nanovoids and microfibrils. Several groups have examined the effects of fillers and matrixes on the thermal analysis of p-aramid containing fabrics and composites, studying thermal and or hydrolytic stability by TGA. Li et al. (2011) and Arrieta et al. (2011) have utilized TGA to assess the thermal stability of p-aramid containing fabrics. Maity et al. (2008) have investigated the effect of fluoronation on the thermal stability of p-aramid fibers. DMA was employed by Li et al. in a 2013 study of the effects of hydrothermal aging on para-aramid fibers. Fig. 19.6 shows the changes noted in storage modulus and tan δ after aging. The authors conclude that the changes noted were caused by chain scission in the noncrystalline regions coupled with rearrangements in intrachain H bonding.

332

Thermal Analysis of Textiles and Fibers

100

150

200

As-spun fiber

–0.075 DSC heat flow (W/g)

Temperature (ºC) 300 350 250

–0.100

Heat treated fiber

–0.125 –0.150

Endo

–0.175

Figure 19.7 DSC of as-spun and heat-treated thermotropic LCP fiber (HBA/HNA 73/2) (Menczel et al., 1997). TA Instrument 2100
2910 DSC, heating rate 10 C/min.

0.90 0.80

A 900 annealed single filament Storage modulus×1011 pa

0.70 0.60 0.50 0.40

tan δ/×10

0.30 0.20 –50 –25

0

25

50

75 100 125 150

1.10 1.00 0.90 0.80 0.70 0.60 0.50 0.40 0.30 0.20

tan δ/×10–1

Modulus/×1011 pa

(A)

Modulus/×1011 pa

1.10 1.00 0.90 0.80 0.70 0.60 0.50 0.40 0.30 0.20 –50 –25

A 900 annealed single filament Storage modulus×1011 pa

tan δ/×10–1 0

25

50

75 100 125 150

1.10 1.00 0.90 0.80 0.70 0.60 0.50 0.40 0.30 0.20

tan δ/×10–1

Temperature (ºC) (B)

Temperature (ºC)

Figure 19.8 DMA of as-spun and heat-treated thermotropic LCP fiber (HBA/HNA 73/27) (Menczel et al., 1997). Perkin
Elmer DMA7, 2 C/min heating. (A) as-spun single filament, (B) annealed single filament.

While there is a rich literature focused on the thermal analysis of thermotropic polyesters (see works of Jaffe, 1981; Jaffe et al., 1997; Menczel et al., 1997), and Romo-Uribe (see, for example, Reyes-Mayer, 2013), and a few papers dealing with other chemistries, there is little recent publication focused on thermotropic liquid

Thermal analysis of liquid crystalline polymers

333

crystalline fibers. Fig. 19.7 shows the DSC curves of as-spun and heat-treated Vectran fibers (Menczel et al., 1997) and Fig. 19.8 shows DMA data for similar samples. The sharpening of the DSC trace with annealing suggests increased perfection in crystalline order (X!N transition), while the DMA results suggest increases in conformational order. In a more recent study of thermally induced reorganizations in thermotropic LCP fibers (Menczel et al., 1997), the authors conclude that the tightening of entangled LCP molecular chains may be a viable mechanism contributing to the properties noted in LCP fibers after annealing close to the melting temperature. Surprisingly, there are a few other papers dealing with the thermal analysis of thermotropic fibers in the literature. Given the potential commercial importance of these fibers, the reader is urged to survey the thermotropic LCP literature [see, e.g., the recent chapter by Jaffe et al. on liquid crystalline polymers in the Encyclopedia of Polymer Science and Technology (2014, revision to be published)], and review LCP compositions being explored by Dingemans (Knijnenberg et al., 2006), and the exploration of potential LCPs based on renewable resource monomers, such as isosorbide (Feng et al., 2010) or 2,5-furandicarboxylic acid (Bhattacharjee et al., 2014). In summary, while LCP fibers have become significant commercial materials for a variety of high-performance applications, the relative difficulty of procuring samples for study in universities and reticence of applied laboratories to publish has served to make the study of these important fibrous materials difficult to systematically study.

References Arrieta, et al., 2011. Polym. Compos. 32 (3), 362
367. Baltussen, J.J., Northolt, M.G., 1996. Polym. Bull. 36, 125
131. Bhattacharjee, Debkumar; Sehanobish, 2014. Kalyan From PCT Int. Appl. WO 2014100256 A2 20140626. Bourbigot, S., Flambard, X., 2002. Fire Mater. 26 (4
5), 155
168. Brown, J.R., Ennis, B.C., 1977. Text. Res. J. 47, 62. Cai, X., Yu, W., 2009. Ind. Text. 62 (4), 187
191. Calundann, G., Jaffe, M., 1982. Proceedings of the Robert A. Welsh Conferences on Chemical Research XXXVI Synthetic Polymers. Calundann, G., Jaffe, M., Yoon, H., 1988. In: Bunsell, A. (Ed.), Fibre Reinforcements for Composite Materials”. Elsevier, Chapter 9. Calundann, G.W., David, H.L., Gorman, F.J., Mininni, R.M., 1976. US Patent 4083829A. Collins, G., Long, B., 1994. J. Appl. Polym. Sci. 53 (5), 587
608. Comet, et al., 2009. Mater. Lett. 63, 279
281. Druzhinina, et al., 2012. Khim. Tekhnol. 13 (6), 345
353. Feng, et al., 2010. ACS Symp. Ser. 10, 3
27. Guth, E., Mark, H., 1934. Monatsh. Chem. 65, 93. Hoegger, E.F., Schaefgen, J.R., Stephens, C.W., 1971. US Patent 3,575,933. Jackson, W.J., 1982. Contemporary Topics in Polymer Science. Springer, Boston, MA.

334

Thermal Analysis of Textiles and Fibers

Jaffe, M., 1981. Fibers. In: Turi, E. (Ed.), Thermal Characterization of Polymer Materials. Academic Press, New York. Jaffe, M., 1987. High modulus polymers, second ed. The Encyclopedia of Polymer Science and Engineering, 7. John Wiley & Sons, United Kingdom, p. 699. Jaffe, M., 1991. J. Stat. Phys. 62 (5/6), 985. Jaffe, M., Jones, R.S., 1985. High performance aramid fibers. In: Lewin, M., Preston, J. (Eds.), Handbook of Fibers: Science and Technology, vol. III. Marcel Dekker. Jaffe, M., Chung, Tai-chi, Friedzon, Y., 2018. Liquid crystalline polymers, Encyclopedia of Polymer Science and Technology, second ed. Wiley-Interscience. Jaffe, M., Menczel, J.D., Bessey, W.E., 1997. Fibers. In: Turi, E. (Ed.), Thermal Characterization of Polymer Materials, second ed. Academic Press, San Diego, CA. Kavesh, S., et al., 2006. US Patent US20070231572A1 (Issue Date 2010, Priority Date 2006). Knijnenberg, A., et al., 2006. Macromolecules 39 (20), 6936. Kwolek, Stephaie, P.W. Morgan, 1966. US Patent 3,287,323. Lammers, M., Klop, E.A., Northolt, M.G., Sikkema, D.J., 1998. Polymer 39 (24), 5999
6005. Liu, X., Yu, Y., 2006. J. Appl. Polym, Sci. 90 (3), 937
944. Li, G., et al., 2011. Huagong Xinxing Cailiao 39 (8), 51
53. Li, C.-S., et al., 2013. J. Appl. Polym. Sci. 128 (2), 1291
1296. Maity, J., Jacob, C., Das, C., Singh, R., 2008. J. Appl. Polym. Sci. 107 (6), 3739
3749. Mark, H., 1936. Pharm. Monatsh. 17, 233. Menczel, J.D., Prime, R.B., 2009. Thermal Analysis of Polymers. Wiley. Menczel, J.D., Collins, G., Saw, K., 1997. J. Therm. Anal. 49, 201. Penn, L., Larsen, F., 1979. J. Appl. Polym. Sci. 23, 59. Perepelkin, et al., 2005. Fibre Chem. 37 (5), 346
351. Reyes-Mayer, Adriana; Alvarado-Tenorio, Bonifacio; Romo-Uribe, Angel; Jaffe, Michael, 2013. From Polymers for Advanced Technologies, 24 (12), 1029
1039. Richardson, I.D., Ward, Ian, 1970. J. Phys. D: Appl. Phys. 3 (5), 643
648. Sawyer, L.C., Jaffe, M., 1986. J. Mater. Sci. 21, 1897. Smith, P., Lemstra, P., 1980. Polymer 21, 1341. Yang, M., et al., 2011. Guangpuxue Ya Guangpu Fensi, ACS Nano. 5 (9), 6945
6954.

Further reading Ciferri, A., Krigbaum, R.W., 2012. In: Meyer, B.R. (Ed.), Polymer Liquid Crystal. Academic Press. Collyer, A.A., 1992. Liquid Crystal Polymers: From Structures to Applications. Elsevier Science Publisher, London. Gupta, V.B., Kothari, V.K., 1997. Manufactured Fibre Technology. Springer Science. Hongu, T., 1997. New Fibers. Woodhead Publishing, Cambridge. Picken, S.J., Sikkema, D.J., Boerstoel, H., Dingermans, T.J., Zwaag, S., 2011. Liq. Cryst. 33 (11
12), 1591
1605. Saw, C.K., Collins, G., Menczel, J., Jaffe, M., 2008. J. Therm. Anal. Calorim. 93 (1), 175
182. Wisss, R.A., Huh, W., Niclais, L., 1987. Polym. Eng. Sci. 27 (9), 684
691.

Thermal analysis of temperature responsive fibrous materials

20

Danmei Sun1, Kashif Iqbal2 and Muhammad Owais Raza Siddiqui3 1 School of Textiles and Design, Heriot-Watt University, United Kingdom, 2Department of Textile processing, National Textile University, Faisalabad, Pakistan, 3Department of Textile Engineering, NED University of Engineering & Technology, Karachi, Pakistan

Abstract This chapter will provide an introduction to the importance of thermal properties of textiles related to physiological comfort, the emerging area of smart textiles, and the benefits of phase-change material (PCM), which would bring for the thermal management solution of textile and clothing system. The working principle of PCM will be described, and the types of PCMs will be reported in details. Fibers incorporated with PCMs and their thermal management properties will be discussed based on the state-of-the-art research studies. The applications of PCMs and the fibers incorporated with PCMs will also be reported.

20.1

Introduction

Thermal properties of textiles play a very important role in determining the comfort of textile and clothing system. Comfort may be divided into four categories containing thermal or thermo-physiological comfort, sensorial comfort, garment fit, and psychological comfort, but the most important one is thermo-physiological comfort (Bhatkhande, 2011). Normal human body temperature is 37 C, which increases up to 38 C, 39 C, or occasionally 40 C during exercise (Havenith et al., 2008). The most comfortable skin temperature is 33.4 C, and when it changes more than 4.5 C below or above the comfort temperature, human body feels uncomfortable (Meng and Hu, 2008). One of the primary functions of clothes is to prevent the temperature of skin to fall or raise too far from the stated comfort temperatures (Mattila, 2006). As far as clothing and comfort is concerned, physical, physiological, and some psychological factors influence the comfort, which is a state of mind. Complete comfort is obtained with the control of thermal and moisture buffering, because the body feels sensation of heat or cold and skin wetness. So, thermal buffering is obtained by using phase-change material (PCM), which, in certain cases, can increase evaporative resistance impairing moisture buffering (Havenith, 2002). Smart textile is an emerging area in textile field, which is becoming more significant by the demand of society through consumer needs. Despite the increasing impact of science and technology, smart textile demands the advancement through interdisciplinary Thermal Analysis of Textiles and Fibers. DOI: https://doi.org/10.1016/B978-0-08-100572-9.00020-3 © 2020 Elsevier Ltd. All rights reserved.

336

Thermal Analysis of Textiles and Fibers

support, such as fashion, design, engineering, technology, human, and life sciences. In textile sector, smart textiles have its vast application in interior textiles, technical textiles, and clothing, in which the last one contains higher percentage in terms of usage of smart textiles (Schwarz et al., 2010). PCMs are kind of smart materials, which were used in clothing by the US National Aeronautics and Space Administration (NASA) in 1980 to make thermoregulated garment for space and to protect apparatus in space with drastic temperature change (Jiang et al., 2008; Cheng et al., 2012; Cox, 1998). This technology was then adopted by Outlast Technologies, based in Boulder, Colorado, which used PCM in textile fibers and fabric coating (Nelson, 2002). Thermal management solution has always been a challenge where high temperature and heat is involved, such as steel mills, glass factories, and sports activities (Yang et al., 2007). To overcome these problems, researchers have been working on smart textiles or intelligent textiles. PCMs are active smart materials, which can sense and react to the environmental stimuli by storing and releasing heat energy (Onofrei et al., 2010). PCMs cannot be applied directly on textiles, because they melt by absorbing heat or crystallize by releasing the same heat. To prevent PCMs from leakage, a process called microencapsulation or nanoencapsulation is used for durability and better application (Mondal, 2008). PCMs are an attractive technique of storing energy in all the available heat energy storage techniques due to high density, compact storage system, and high latent heat (Mondal, 2008; Farid et al., 2004). The pioneer study of PCM was applied for spacecraft on small scale and then on large scale in buildings and solar energy systems to build thermal energy storage system (Buddhi et al., 1987; Lacroix, 1993; Pauken et al., 2007). A large number of inorganic and organic PCMs are available in the temperature range of 25 C to 190 C (Agbossou et al., 2010; Alkan, 2006; Choi et al., 2001; Tyagi and Buddhi, 2007; Zalba et al., 2003). The organic PCMs ranging from 18 C to 65 C are used in textiles and buildings to enhance thermal comfort effect. Among all the PCMs, n-octadecane is usually used for the textile application having melting point of 28 C (Feldman et al., 1986). The developers claim that incorporation of PCMs in textiles will perform buffering effect keeping the skin temperature constant against extreme weather, hence, prolonging thermal comfort for the wearer. They also claim that using PCMs can decrease the fabric thickness required to protect the human body from cold environment (Pause, 1995; Cox, 1998). Currently, PCMs are being used in different textiles, including bedding, apparel, footwear, and nonwoven under the trade names of Outlast, TemperTex, and ComforTemp. Outlast Technologies has succeeded in marketing viscose and PAN (polyacrylonitrile) fibers incorporated with microPCM (Hartmann et al., 2004).

20.2

Working principle of phase-change material

PCMs store energy as latent heat, which is the most suitable way of storing high amount of energy with a smaller difference of temperature. This high amount of heat is released to surroundings during reverse cooling process called crystallization. By incorporating PCMs to textiles, their heat-storage capacity can be substantially enhanced (Farid et al., 2004).

Temperature

Thermal analysis of temperature responsive fibrous materials

337

Phase-change temperature

Temperature control range

Stored heat

Figure 20.1 Working principle of PCM.

During melting and crystallization, the temperature of PCM does not change significantly. The increase in temperature of surroundings causes PCM to melt by absorbing heat, and when temperature of surrounding environment falls down, PCM gets solidified releasing the same amount of heat energy to environment. This interesting phenomenon makes it very important in practical applications (Farid et al., 2004). The working principle of PCM is demonstrated in Fig. 20.1, where there is a temperature control range near to the phase-change temperature. The period of temperature will depend on the amount of latent heat (J) stored in the material. Chemical bonding is responsible for the previous phenomenon as increase in temperature goes to break chemical bonding in the molecules of PCM material causes material to melt resulting in storing heat energy, which subsequently released during crystallization process restoring chemical bond (Bendkowska and Wrzosek, 2009). To control the phase change between liquid and solid of PCMs, a technique is repeatedly used to convert the liquid state PCM into solid, which is called encapsulation. Microencapsulation is the process of manufacturing polymeric particles in the range from nanometers to millimeters, in which two main morphologies exist as microencapsulation and microspheres. Microsphere is the process, in which active substance is enclosed within the polymeric network, while microencapsulation exhibits the reservoir structure in which the active substance is surrounded by a polymeric wall called capsule shell, shown in Fig. 20.2. The methods of microencapsulation can be classified into three groups: mechanical method, physicochemical method, and chemical method.

20.3

Types of phase-change material

PCMs are theoretically able to change their phase nearly constant temperature, and therefore they are able to store large amount of energy (Ku¨rklu¨, 1997). More than

338

Thermal Analysis of Textiles and Fibers

Polymer shell

PCM core

Figure 20.2 Structure of capsulated PCM.

500 natural and synthetic PCMs are known in addition to water, but they differ in melting points and heat-storage capacities (Pause, 2002). Among all used PCMs for textiles, the most common is paraffin that can be microencapsulated and then either incorporated into fiber or applied via coating on textiles. Some types of PCM with phase-change temperature close to human skin temperature are described in the following subsections.

20.3.1 Hydrated inorganic salt Inorganic salts containing n water molecules can be used as PCM in thermoregulating textiles, which usually have phase-change temperature in the range of 20 C 40 C. Glauber’s salt (Na2SO4
10H2O) is very attractive and convenient because of its physical and chemical properties. It has a melting temperature of 32.4 C that is very suitable for textiles and has a large amount of latent heat of 254 J/g (Canbazo˘glu et al., 2005). In the thermal energy system using Glauber’s salt as a PCM, the thermal energy is produced by the chemical reaction between decahydrate crystal and water solution (Saito et al., 2001). Hydrated salts are found useful in thermal energy storage system due to their high volumetric storage capacity, relative high thermal conductivity, and lower cost as compared to paraffin. Glauber’s salt containing 44% of sodium sulfate and 56% of water by weight was studied and investigated by Biswas (1977) and Marks (1980). Li et al. successfully made CaCl
6H2O- and MgCl2
6H2O-based PCM. The phase-change temperature and latent heat were measured by differential scanning calorimetry (DSC) as shown in Fig. 20.3 (Zhang et al., 2005). The phase-change temperature is 21.4 C, and the latent is 102.3 J/g.

20.3.2 Organic hydrocarbons Long-chain hydrocarbons are the by-product of oil refining containing general formula of CnH2n12. They are suitable for various applications as they have wide range of melting temperature depending upon the number of carbon atoms. They are usually nontoxic, inexpensive and have large source of raw materials. By selecting the number of carbon atoms present in hydrocarbons, the required phase-change temperature can be tailored for specific application. Sari et al. (2009) reported their

Thermal analysis of temperature responsive fibrous materials

0

339

Onset: 21.4ºC Latent heat: 102.3 J/g

Heat flow (mW/mg)

–10

–20

–30 –40 –50 0

10

20

30

40

50

60

70

80

Temperature (ºC)

Figure 20.3 DSC melting curve of CaCl2
H2O 1 25 wt.% MgCl2
6H2O 1 3 wt.% SrCl2
6H2O 1 1 wt.% SrCO3 1 0.5 wt.% hydroxyethyl cellulose (HEC) (Zhang et al., 2005).

Figure 20.4 SEM image of PMMA/octacosane microcapsules (Sari et al., 2009).

research work with preparation and characterization of polymethylmetracrylate (PMMA) microcapsules containing n-octacosane as PCM for thermal energy storage. The surface morphology, particle size, and particle size distribution were studied by scanning electron microscopy (SEM) presented in Fig. 20.4. Thermal

340

Thermal Analysis of Textiles and Fibers

properties and thermal stability of microencapsulated octacosane were determined using DSC and thermal gravimetric analysis (TGA). The melting and freezing temperatures and the latent heats of the microencapsulated octacosane as PCM were measured as 50.6 C and 53.2 C, 86.4 and 88.5 J/g, respectively, through DSC analysis. The chemical stability of PMMA/octacosane microcapsules after repeated thermal cycling was also investigated by fourier transform infrared spectroscopy (FT-IR) analysis as shown in Fig. 20.5. The chemical structure of microcapsules was not affected by thermal cycling reflected by the consistency of peaks in the samples before and after thermal cycling. These are agreed well with research work carried out by Uddin et al. (2002). In the meantime the thermal reliability of the PMMA/octacosane microcapsules was proved in terms of the change in its phasechange temperature after repeated 5000 thermal cycling process. n-Octadecane and eicosane containing carbon atoms 18 and 20, respectively, are suitable for textiles, because their phase-change temperature lies near the human

PMMA/Octacosane

PMMA/Octacosane after 5000 thermal cycling

%T

4000

3000

2000

1000

400

–1)

Wavenumber (cm

Figure 20.5 FT-IR spectra for PMMA/octacosane microcapsules before and after thermal cycling (Sari et al., 2009).

Thermal analysis of temperature responsive fibrous materials

341

skin temperature (Zuckerman et al., 2003). The phase-change temperature and latent heat of organic hydrocarbon are 28 C 30 C and 244.8 J/g, respectively.

20.3.3 Polyethylene glycol Polyethylene glycol (PEG) is another important type of PCM for textile applications. The repeating unit in PEG is oxyethylene ( O CH2 CH2 ) containing hydroxyl group on either side of the chain. The melting point of PEG depends on its molecular weight and is proportional as the molecular weight increases. The phase-change temperature of PEG can be determined using DSC (Pielichowski and Flejtuch, 2002). PEG with degree of polymerization 1000 has phase-change temperature of 35 C, while PEG with degree of polymerization 20,000 has melting temperature of 63 C (Craig and Newton, 1991; Hopp et al., 2000). Jiang et al. (2016) synthesized a dual-functional magnetic microcapsules containing a PCM core and an organo-silica shell for the electromagnetic shielding and thermal regulating applications. Fig. 20.6 shows the resulting DSC curves where

Figure 20.6 TGA and difference thermogravimetry (DTG) curves of (A) pure PEG and the magnetic microcapsules synthesized at organosilanes/PEG weight ratios of (B) 75/25, (C) 67/ 33, and (D) 50/50 (Jiang et al., 2016).

342

Thermal Analysis of Textiles and Fibers

the areas under the peaks indicate the amount of latent heat contained using different organosilanes/PEG weight ratios.

20.4

Phase-change material incorporated fibers and their thermal management performance

Both PCM containing yarns and fabrics can be produced through various manufacturing techniques. There are three ways in making thermoregulating filament containing PCMs: (1) through composite spinning process in which PCM is directly mixed with melt spun or solvent spun polymers as core/sheath filament were PCM as core (Salyer, 1999); (2) through hollow fiber spinning method, in which hollow fiber is filled with PCM (Vigo et al., 1990); and (3) either melt spun or solvent spun incorporating encapsulated PCM, and provides the PCM textiles with good chemical and thermal storage stability (Zhang et al., 2005). There are three approaches in producing thermos-regulating fabrics: (1) producing PCM fabrics using PCM-incorporated yarns; (2) applying microencapsulated PCMs onto fabrics via coating or pad-dry-cure process with the help of binder to ensure the binding of capsules with fabric along with some ancillary chemicals. Different processes of coating, such as knife over air, knife over roll, dip coating, and transfer coating, can be used to apply polymer matrix on textiles (Nelson, 2002; Zuckerman et al., 2003); and (3) applying microencapsulated PCM onto fabrics through lamination process. Microencapsulated PCMs can be incorporated into fiber through conventional filament spinning process. The microencapsulated PCMs are first added to the fiber forming solution and extruded with conventional solvent spinning or can be mixed with polymer prior to melt spinning to get thermoregulating melt spun filament yarn. Most of the work related to the incorporation of microPCM into fiber is patented, and some literature is published about application via coating or pad-dry-cure techniques. Many researchers have tried to incorporate capsules containing PCM into man-made fibers. Bryant and Colvin (1988) succeeded in incorporating capsules containing eicosane as PCM into viscose rayon and acrylic fiber by wet or dry spinning method and patented their work. These filament yarns gave thermoregulating characteristics when subjected to heat and cold. Viscose rayon and acrylic successfully came into the market containing leak resistant capsules by an American manufacturer, Outlast. Outlast Technologies assigned many projects to researchers to get fiber enhanced with reversible thermal properties. Hartmann and Magill (2004) developed viscose rayon fiber containing microencapsulated PCMs and patented in 2010, which was the continuation of their previous work. The latent heat mentioned in their patent was from 1 20 J/g in dependent of the different embodiments. Zhang et al. (2006) prepared flame-retardant PAN-vinylidene chloride fiber through wet spinning containing n-octadecane as PCM and urea melamine formaldehyde copolymer as shell. The incorporated amount of MicroPCM (MPCM) studied was 0% 40%, which explored the smooth spinning up to 30%. The enthalpy of

Thermal analysis of temperature responsive fibrous materials

343

crystallization was determined by DSC from 220 C to 320 C with temperature gradient of 10 C/min under nitrogen atmosphere showing 30 and 44 J/g for fiber containing 30% and 40% of microPCM, respectively. SEM result showed that the fibers that contain microPCM have rough surface, more density, and cross-sectional holes as compared to non-microPCM. Incorporation of encapsulated PCM into melt spun fiber is very problematic due to high temperature (200 C 380 C) and high pressure (3000 psi) used during the process. At such high temperature and pressure, normally, PCMs undergo thermally induced decomposition and oxidation. The factors affecting decomposition are temperature, pressure, and duration of time during which PCMs are exposed. This thermally induced decomposition leads to the formation of lower molecular products or isomers, which not only reduce thermal regulatory properties but also impair melt spinning process (Hartmann, 2004). Due to the thermal and shear stability, M/F capsules were used, which cause some environmental issues during the preparation of capsules. Frank and Biastoch (2001) published patent in 2001 about preparing M/F capsules through condensation process of melamine and formaldehyde and added urea as formaldehyde scavenger prior to curing of resin. Later on, Li et al. (2007) manufactured microPCM using melamine formaldehyde shell by putting aluminum chloride during the microencapsulation process to reduce the residual formaldehyde content. Outlast Technologies is anxiously waiting for melt spun thermoregulating fiber. Bryant (1999) in his paper mentioned the problem related to melt spinning filament yarns containing microPCM and suggested that the lower particle size preferably less than 10 µm is suitable for melt spinning filament yarns. Hartmann (2004) prepared stabilized PCMs by mixing PCMs with stabilizing agents consisting of thermal stabilizers and antioxidants that can be used in synthetic fiber manufacturing. The encapsulation of PCM was done by incorporating PCM into hollow capsules containing alone PCM or mixture of PCM and stabilizers in the internal cavity and stabilizing agents in outer cavity. Hartmann et al. (2004) in another patent mentioned the development of melt spinnable concentrate pellets containing PCMs either in microencapsulated form or nonencapsulated form absorbed into the carrier polymer. The polymer matrix was composed of either one type of thermoplastic polymer or a combination of them. They claimed that these concrete pellets can be used with other thermoplastic polymers to form extrusion products, including monofilament fiber, but there was no filament made with such pellets. Magill et al. (2005) prepared a multicomponent fiber containing PCM by keeping the PCM in core and any thermoplastic or elastic fiber in the outer sheath using polyester, nylon, and many other polymers. They claimed in their patent that material contains 6.9 and 8.4 J/g of latent heat in different embodiments of core/sheath yarn. Gao et al. (2009) developed PAN fiber containing MA (methyl acrylate) copolymer by melt spinning technique. The fiber was spun at 200 C containing 5% 25% of microPCM. The enthalpies of crystallization of fibers containing 20% and 25%

344

Thermal Analysis of Textiles and Fibers

of PCM were 21 and 25 J/g, respectively, tested by DSC. The SEM analysis showed that the fibers containing PCM are coarser, while non-microPCM fibers are smooth and the fibers containing capsules have microporous and defective structure as shown in Fig. 20.7. Wen et al. (2015) developed a core sheath phase-change microfiber with paraffin in the core and polyvinyl butyral (PVB) polymer being the shell. The setup is shown in Fig. 20.8. The microstructures, thermal characteristics, and mechanical

Figure 20.7 SEM images of PAN/MA fiber (A) with PCM and (B) without PCM (Gao et al., 2009). Inner fluid (Qi) Injection tube

Middle fluid (Qm)

Outer fluid (Qo) Transition tube

Collection tube

Figure 20.8 The setup of coaxial micro device and formation process of core sheath phasechange microfibers (Wen et al., 2015).

Thermal analysis of temperature responsive fibrous materials

345

strengths of microfibers were studied systematically by SEM, DSC, TGA, and a tensile tester. The enthalpies of microfibers increased with the increase of PCM content in microfibers. The maximum melting enthalpy and crystallization enthalpy of microfibers were found to be about 128.2 and 124.0 J/g, respectively, and the corresponding encapsulation ratio was as high as 70%. The prepared phase-change microfibers exhibited stable and satisfactory thermal characteristics especially thermoregulating property. More recently, Iqbal and Sun (2014) reported a novel type of smart monofilament fiber polypropylene (PP) incorporated with microencapsulated PCM through melt spinning process. Up to 12% of PCM microcapsules are successfully incorporated into the PP monofilament showing 9.2 J/g of latent heat. Fig. 20.9 shows the SEM images of PP filament incorporated with microencapsulated PCM as comparison to that of without. The PCM-incorporated filament shows rough surface and many particles present in the surface of the filament (A), indicating the presence of the microencapsulated PCM in comparison to the fiber without PCM (B). The peaks in DSC graph shown in Fig. 20.10 indicate the latent heat in filament yarns containing different amount of microencapsulated PCM compared to the filament yarn without PCM, and it can be seen that no peak is appeared throughout the temperature range. The areas under the peaks indicate the latent heat stored in the filament yarn containing microencapsulated PCM. The peak of curve increases with the increase of the amount of PCM in the filament yarn, indicating the increase in the latent heat with the increase in the amount of PCM. Furthermore, Sun and Iqbal (2015) investigated the thermal regulating effect of a 70% core and 30% sheath yarn with PCM in the core using finite element method. Fig. 20.11 shows the finite element core/sheath yarn model: (A) showing the two parts of the yarn where PCM n-octadecane and sheath are composed of PP and (B) the heat-transfer effect of the yarn model after postprocessing calculation. The model was validated by DSC experimental results.

Figure 20.9 PP fiber incorporated with microencapsulated PCM (A) with PCM, and (B) without PCM.

346

Thermal Analysis of Textiles and Fibers

3500 PP

Endothermic

3000

3% MPCM

2500

12% MPCM

2000 1500 1000 500 25

30

35

40

Temperature (ºC)

Figure 20.10 DSC graph showing latent heat of filaments containing microencapsulated PCM.

Figure 20.11 Finite element analysis of core (PCM)/sheath yarn.

The microPCM in the yarn provides thermal protection through storing energy in the form of latent heat and does not allow the temperature to rise for certain period of time. This temperature delay with respect to time can be predicted using ABAQUS postprocessing, which cannot be tested experimentally. The thermal regulating effect of the core sheath yarn was analyzed and comparison made to the yarn without PCM shown in Fig. 20.12. The yarn without microPCM reaches the highest temperature very quickly, while the yarn containing PCM in the core is delayed to reach the same temperature level. The dashed box in Fig. 20.12 shows the thermoregulation zone of the core sheath yarn containing large amount of PCM.

Thermal analysis of temperature responsive fibrous materials

347

40 38

Temperature (ºC)

36 34 32 Yarn without PCM

30 Core (PCM)/Sheath yarn

28 26 24

Temperature regulation zone

22 20 0

10

20 Time (min)

30

40

Figure 20.12 Thermal regulation effect versus time.

The thermoregulation zone is also called the thermal arrest zone, which means that the temperature is arrested for some time because of the phase-change effect. In Fig. 20.12 the curve of yarn without microPCM proceeds linearly over time showing no temperature arrest, whereas the yarn containing microPCM shows thermal arrest, which lasts as long as the PCM changes its phase completely from solid to liquid. After the latent heat effect completes, the temperature again starts rising and leaves the thermoregulation zone. The heat transfer inside the cross section of simulated yarn with and without PCM at the same time interval from finite element model analysis is shown in Fig. 20.13. The yarn without PCM reaches the temperature of 31 C 32 C, whereas the core/sheath yarn containing PCM in core reaches to 26 C after 2000 seconds, and the larger blue region indicates better thermoregulation effect. Therefore the simulated core/sheath yarn provides better thermoregulation effect.

20.5

Applications of phase-change material and its incorporated fibers

Zhu et al. (2015) studied the thermal protective performance of protective fabrics for firefighters by incorporating PCM. The protective fabrics were composed of four layers with three fabric layers incorporated with PCMs apart from the outer shell layer, shown in Fig. 20.14. Three types of commercial PCM materials with melting temperature range of 29 C 35 C, 47 C 53 C, and 77 C 85 C were used in their study. It was found that all types of composite fabrics containing PCMs had a higher heat protection performance than that of the fabrics without PCM.

348

Thermal Analysis of Textiles and Fibers

Figure 20.13 Contour of thermoregulation effect in yarns.

Figure 20.14 PCM configuration within firefighters’ protective fabrics (FFPF) specimen assembly (Zhu et al., 2015).

Bartkowiak and Dbrowska (2012) reported that two types of knitted fabric with fibers incorporated with PCMs were designed and manufactured for underbarrier protective clothing. The biophysical properties, such as thermal resistance, water vapor resistance, air permeability, and hygroscopicity, and thermoregulation properties, such as enthalpy, thermal recovery facility (TRF), and sorption of the fabrics

Thermal analysis of temperature responsive fibrous materials

349

Figure 20.15 Finite element geometric model (Iqbal et al., 2015).

34 Temperature (ºC)

32 30 28 26

Fabric without MPCM Fabric with 10% MPCM

24 22 20 0

200

400

600

800

1000 1200 1400 1600 1800 2000 Time (s)

Figure 20.16 Thermoregulating effect versus time (Iqbal et al., 2015).

were studied. The impact of the clothing designed containing PCMs, on the underwear microclimate, with an indication of their ability of taking off excess heat was assessed. Iqbal et al. (2015) developed finite element model based on the actual woven fabric made of yarn incorporated with paraffin-based microcapsules. Its heattransfer property was investigated based on the validated model after postprocessing calculation. The delay in temperature rises as a function of time is also predicted, which is not possible to be determined through experiment. The geometric model parts are shown in Fig. 20.15, which consist of yarn, microPCMs, and air. The magnified and cross-sectional image of microPCM has been shown in Fig. 20.15. The green part shows the core part, which comprises PCM, while the encapsulating shell part has been shown in dark peach color. Fig. 20.16 shows the time-dependent temperature changes for the fabrics with and without PCM. Overall, the temperature increases as the time increases.

350

Thermal Analysis of Textiles and Fibers

However, for the PCM containing fabrics, near the melting temperature of PCM, the graph proceeds linearly with time allows storing latent heat and resulting in the delay in the increase of temperature.

20.6

Conclusions and future research

In this chapter the working principle and thermal management behavior of PCM, types of PCMs, fibers, and their thermal-related properties, as well as applications of PCMs, fibers containing PCMs, and their thermal properties have been discussed. Future research on thermal responsive fibers and textiles through incorporating PCMs may consider the following aspects: 1. Design and manufacture of other PCM-incorporated melt spun fibers, such as polyester and nylon, due to the versatile applications of the two fiber materials. 2. The development of nanoscaled PCM capsules for microfibers containing PCMs. 3. The development of micro- or nanoscaled capsules using other PCMs, such as Glauber’s salt, which has phase-change temperature closer to skin comfort temperature for thermal responsive and comfort textiles and clothing.

References Agbossou, A., Zhang, Q., Sebald, G., Guyomar, D., 2010. Solar micro-energy harvesting based on thermoelectric and latent heat effects. Part I: theoretical analysis. Sens. Actuators A: Phys. 163, 277 283. Alkan, C., 2006. Enthalpy of melting and solidification of sulfonated paraffins as PCMs for thermal energy storage. Thermochim. Acta 451, 126 130. Bartkowiak, G., Dbrowska, A., 2012. Assessment of the thermoregulation properties of textiles with fibres containing PCMs on the basis of laboratory experiments. Fibres Text. East. Eur. 20 (1 (90)), 47 52. Bendkowska, W., Wrzosek, H., 2009. Experimental study of the thermoregulating properties of nonwovens treated with microencapsulated PCM. Fibres Text. East. Eur. 17, 76. Bhatkhande, P.S., 2011. Development of Thermo-Regulating Fabric Using PCM (PCM) (Master’s Theses and Doctoral Dissertations). Eastern Michigan University. Biswas, D.R., 1977. Thermal energy storage using sodium sulfate decahydrate and water. Sol. Energy 19, 99 100. Bryant, Y.G., 1999. Melt spun fibres containing microencapsulated PCM. In: Advances in Heat and Mass Transfer in Biotechnology, Amer Society of Mechanical, HTD— vol. 363/BED—vol. 44, pp. 225 234. ISBN: 0791816435. Bryant, Y., Colvin, D., 1988. Fiber With Reversible Enhanced Thermal Storage Properties and Fabrics Made There From, US Patent 4756958. Buddhi, D., Sawhney, R., Sehgal, P., Bansal, N., 1987. A simplification of the differential thermal analysis method to determine the latent heat of fusion of PCMs. J. Phys. D: Appl. Phys. 20, 1601. ´ .G., Akarsu, F., 2005. Canbazo˘glu, S., Sahinaslan, ¸ A., Ekmekyapar, A., Aksoy, Y Enhancement of solar thermal energy storage performance using sodium thiosulfate pentahydrate of a conventional solar water-heating system. Energy Build. 37, 235 242.

Thermal analysis of temperature responsive fibrous materials

351

Cheng, W.L., Liu, N., Wu, W.F., 2012. Studies on thermal properties and thermal control effectiveness of a new shape-stabilized PCM with high thermal conductivity. Appl. Therm. Eng. 36, 345 352. Choi, J.K., Lee, J.G., Kim, J.H., Yang, H.S., 2001. Preparation of microcapsules containing PCMs as heat transfer media by in-situ polymerization. J. Ind. Eng. Chem. 7, 358 362. Cox, R., 1998. Synopsis of the new thermal regulating fiber outlast. Chem. Fibers Int. 48, 475 476. Craig, D., Newton, J., 1991. Characterisation of polyethylene glycols using differential scanning calorimetry. Int. J. Pharm. 74, 33 41. Farid, M.M., Khudhair, A.M., Razack, S. A.K., Al-Hallaj, S., 2004. A review on phase change energy storage: materials and applications. Energy Convers. Manage. 45, 1597 1615. Feldman, D., Shapiro, M.M., Banu, D., 1986. Organic-PCMs for thermal-energy storage. Sol. Energy Mater. 13, 1 10. Frank, G., Biastoch, R., 2001. Low-Formaldehyde Dispersion of Microcapsules of Melamine-Formaldehyde Resins, Patent 6224795. Gao, X., Han, N., Zhang, X., Yu, W., 2009. Melt-processable acrylonitrile methyl acrylate copolymers and melt-spun fibers containing microPCMs. J. Mater. Sci. 44, 5877 5884. Hartmann, M., Worley, J.B., North, M., 2004. Cellulosic Fibers Having Enhanced Reversible Thermal Properties and Methods of Forming Thereof, US 10/638,290. Hartmann, M.H., 2004. Stable PCMs for Use in Temperature Regulating Synthetic Fibers, Fabrics and Textiles, US Patent 6689466. Hartmann, M.H., Magill, M.C., 2004. Melt Spinable Concentrate Pellets Having Enhanced Reversible Thermal Properties, Patent 6793856 B2. Havenith, G., 2002. The interaction of clothing and thermoregulation. Exog. Dermatol. 1, 221 230. Havenith, G., Smith, C., Fukazawa, T., 2008. The skin interface—meeting point of physiology and clothing science. J. Fiber Bioeng. Inf. 1, 93 98. Hopp, B., Smausz, T., Tomba´cz, E., Wittmann, T., Igna´cz, F., 2000. Solid state and liquid ablation of polyethylene-glycol 1000: temperature dependence. Opt. Commun. 181, 337 343. Iqbal, K., Sun, D., 2014. Development of thermo-regulating polypropylene fibre containing microencapsulated PCMs. Renew. Energy 71, 473 479. Iqbal, K., Sun, D., Stylios, G., Lim, T., Corne, D., 2015. FE analysis of thermal properties of woven fabric constructed by yarn incorporated with microencapsulated PCMs. Fibers Polym. 16 (11), 2497 2503. Jiang, M., Song, X., Xu, J., Ye, G., 2008. Preparation of a new thermal regulating fiber based on PVA and paraffin. Sol. Energy Mater. Sol. Cells 92, 1657 1660. Jiang, F., Wang, X., Wu, D., 2016. Magnetic microencapsulated PCMs with an organo silica shell: design, synthesis and application for electromagnetic shielding and thermal regulating polyimide films. Energy 98, 225 239. Ku¨rklu¨, A., 1997. Thermal performance of a tapered store containing tubes of PCM: cooling cycle. Energy Convers. Manage. 38, 333 340. Lacroix, M., 1993. Study of the heat-transfer behavior of a latent-heat thermal-energy storage unit with a finned tube. Int. J. Heat Mass Transfer 36, 2083 2092. Li, W., Wang, J., Wang, X., Wu, S., Zhang, X., 2007. Effects of ammonium chloride and heat treatment on residual formaldehyde contents of melamine-formaldehyde microcapsules. Colloid Polym. Sci. 285, 1691 1697.

352

Thermal Analysis of Textiles and Fibers

Magill, M.C., Hartmann, M.H., Haggard, J.S., 2005. Multi-Component Fibers Having Enhanced Reversible Thermal Properties and Methods of Manufacturing Thereof, Patent 6855422 B2. Marks, S., 1980. An investigation of the thermal energy storage capacity of Glauber’s salt with respect to thermal cycling. Sol. Energy 25, 255 258. Mattila, H., 2006. Intelligent Textiles and Clothing. Woodhead Publishing, USA. Meng, Q.H., Hu, J.L., 2008. A poly(ethylene glycol)-based smart PCM. Sol. Energy Mater. Sol. Cells 92, 1260 1268. Mondal, S., 2008. PCMs for smart textiles—an overview. Appl. Therm. Eng. 28, 1536 1550. Nelson, G., 2002. Application of microencapsulation in textiles. Int. J. Pharm. 242, 55 62. Onofrei, E., Rocha, A.M., Catarino, A., 2010. Textiles integrating PCMs a review. Buletinul Institutului Politehnic din Ia¸si 2, 99 110. Pauken, M., Emis, N., Watkins, B., 2007. Thermal energy storage technology developments. Space Technol. Appl. Int. Forum—STAIF 2007 880, 412 420. Pause, B., 1995. Development of heat and cold insulating membrane structures with PCM. J. Ind. Text. 25, 59 68. Pause, B., 2002. Driving more comfortably with PCMs. Tech. Text. Int. 11, 24 27. Pielichowski, K., Flejtuch, K., 2002. Differential scanning calorimetry studies on poly (ethylene glycol) with different molecular weights for thermal energy storage materials. Polym. Adv. Technol. 13, 690 696. Saito, A., Okawa, S., Shintani, T., Iwamoto, R., 2001. On the heat removal characteristics and the analytical model of a thermal energy storage capsule using gelled Glauber’s salt as the PCM. Int. J. Heat Mass Transfer 44, 4693 4701. Salyer, I.O., 1999. A Clothing Formed From a Melt Blends Comprising a Polyolefin Selected From the Group Consisting of High Density Polyethylene, Low Density Polyethylene, and a High Melting Polypropylene, an Ethylene-Vinyl Acetate Copolymer, Silica, US Patent 5885475. Sari, A., Alkan, C., Karaipekli, A., Uzun, O., 2009. Microencapsulated n-octacosane as PCM for thermal energy storage. Sol. Energy 83, 1757 1763. Schwarz, A., Van Langenhove, L., Guermonprez, P., Deguillemont, D., 2010. A roadmap on smart textiles. Text. Prog. 42, 99 180. Sun, D., Iqbal, K., 2015. Investigating thermal properties of filament yarn containing PCMs. In: Proceedings of 15th AUTEX Conference, Bucharest, Romania, AUTEX, 10 12 June 2015. Tyagi, V.V., Buddhi, D., 2007. PCM thermal storage in buildings: a state of art. Renew. Sustain. Energy Rev. 11, 1146 1166. Uddin, M.S., Zhu, H.J., Hawlader, M.N.A., 2002. Effects of cyclic operation on the characteristics of a microencapsulated PCM storage material. Int. J. Sol. Energy 22, 105 114. Vigo, T.L., Zimmerman, C.M., Bruno, J.S., Danna, G.F., 1990. Temperature Adaptable Textile Fibers and Method of Preparing Same, US Patent 4908238. Wen, G., Xie, R., He, X., Wang, W., Ju, X., Chu, L., 2015. Microfluidic fabrication and thermal characteristics of core shell phase change microfibers with high paraffin content. Appl. Therm. Eng. 87 (5), 471 480. Yang, X.B., Zheng, J.P., Bai, Y., Tian, F.J., Yuan, J., Sun, J.Y., 2007. Using lymphocyte and plasma Hsp70 as biomarkers for assessing coke oven exposure among steel workers. Environ. Health Perspect. 115, 1573 1577.

Thermal analysis of temperature responsive fibrous materials

353

Zalba, B., Marin, J.M., Cabeza, L.F., Mehling, H., 2003. Review on thermal energy storage with phase change: materials, heat transfer analysis and applications. Appl. Therm. Eng. 23, 251 283. Zhang, X., Wang, X., Tao, X., Yick, K., 2005. Energy storage polymer/microPCMs blended chips and thermo-regulated fibers. J. Mater. Sci. 40, 3729 3734. Zhang, X., Wang, X., Tao, X., Yick, K., 2006. Structures and properties of wet spun thermoregulated polyacrylonitrile-vinylidene chloride fibers. Text. Res. J. 76, 351 359. Zhu, F., Feng, Q., Liu, R., Yu, B., Zhou, Y., 2015. Enhancing the thermal protective performance of firefighters’ protective fabrics by incorporating PCM. Fibres Text. East. Eur. 23 (2), 68 73 (110). Zuckerman, J.L., Pushaw, R.J., Perry, B.T., Wyner, D.M., 2003. Fabric Coating Containing Energy Absorbing PCM and Method of Manufacturing Same, US 6514362 B1.

Further reading Li, G., Zhang, B., Li, X., Zhou, Y., Sun, Q., Yun, Q., 2014. The preparation, characterization and modification of a new phase change material: CaCl2
6H2O MgCl2
6H2O eutectic hydrate salt. Sol. Energy Mater. Sol. Cells 126, 51 55.

Thermal characterization of fire-protective fabrics

21

Sumit Mandal1,*, Sabyasachi Gaan2, Martin Camenzind1, Simon Annaheim1 and Rene´ M. Rossi1 1 Empa, Swiss Federal Laboratories for Materials Science and Technology, Laboratory for Biomimetic Membranes and Textiles, St. Gallen, Switzerland, 2Empa, Swiss Federal Laboratories for Materials Science and Technology, Laboratory for Advanced Fibers, St. Gallen, Switzerland

Abstract This chapter discusses about the thermal characterization of fire protective fabrics. In Section 21.1, the manufacturing process and uses of fire protective fabrics are highlighted. Generally, fire protective fabrics are used as raw materials for the manufacturing of workwear for firefighters, oil and gas industry labors, defense personnel, and cooks/chefs. Indeed, the evaluation of fire protective performance of fabrics is necessary for providing the safe working environments to these workers. Considering this, different standardized and customized test methods are developed for evaluating the fire protective performance of fabrics. In this chapter, these test methods are thoroughly described. By using these test methods, many researchers also evaluated and analyzed the fire protective performance of fabrics in the last few decades. In this chapter, their research is critically reviewed for better understanding of fire protective performance of fabrics. Thereafter, some key issues related to the fire protective fabrics are presented. By resolving these key issues, textile or materials engineers could develop a high-performance fire protective fabric, and this fabric based workwear could provide the better protection and safety to the workers engage in different high-risk sectors such as firefighting, oil and gas industry, defense, and restaurants.

21.1

Introduction

Fire-protective fabrics are generally manufactured from textile fibers that are noncombustible in minimum atmospheric oxygen concentration of 21% (Fig. 21.1) (Bajaj, 1992). These noncombustible fibers can be classified as (1) chemically modified fire-retardant (FR) fibers and (2) inherently fire-resistant fibers (Fig. 21.1) (Kilinc, 2013). Chemically modified FR fibers are originally combustible natural (e.g., cotton, wool), regenerated (e.g., viscose), or synthetic (e.g., polyester, nylon,




Current Affiliation: Department of Design, Housing, and Merchandising, College of Human Sciences, Oklahoma State University, Stillwater, Oklahoma, United States

Thermal Analysis of Textiles and Fibers. DOI: https://doi.org/10.1016/B978-0-08-100572-9.00021-5 © 2020 Elsevier Ltd. All rights reserved.

356

Thermal Analysis of Textiles and Fibers

Textile fibers

Spinning Weaving or knitting

Weaving or knitting

Yarns

Yarns

FR treatment Fire protective woven fabric

(e.g., halogen, nitrogen)

Fabric

FR Treatment Fire protective nonwoven fabric

(e.g., halogen, nitrogen)

Nonwoven process

Nonwoven process

Combustible fibers 1. Natural/regenerated fibers (e.g., cotton, viscose) 2. Synthetic fibers (e.g., polyester, nylon)

(Chemical, mechanical)

(chemical, mechanical)

FR treatment (e.g., halogen, nitrogen)

Spinning

Noncombustible fibers 1. Chemically modified fire-retardant fibers (e.g., FR cotton, FR polyester); 2. Inherently fire-resistant fibers (e.g., aramid, polyimide)

Nonwoven fabric

Figure 21.1 Manufacturing processes of fire-protective fabrics.

acrylic) fibers; later, these substrate fibers are finished/coated with FR chemical substances (halogen, nitrogen, silicon, and phosphorous) to make them noncombustible (e.g., FR cotton, FR polyester) (Gaan et al., 2011a,b; Liang et al., 2013; Neisius et al., 2014). Notably, these chemically finished fibers can lose their fire retardancy as result of wear, age, or frequent washing. For synthetic substrate fibers (e.g., polyester) with permanent fire retardancy, the flame-retardant substances are sometimes incorporated into the polymerization during the fiber manufacturing process through melt spinning or doped in the spinning bath during the production of fibers by solution spinning (e.g., Trevira CS) (Fabric for Ideas, n.d.). Inherently, fire-resistant fibers are noncombustible synthetic fibers (e.g., aramid, polyimide, polybenzimidazole) that are generally manufactured through the melt or solution spinning of thermoplastic polymers. After producing the chemically modified FR or inherently fire-resistant fibers, these fibers are used as raw materials in the spinning process to manufacture yarns; subsequently, these yarns are used as raw materials (warp and weft) in the weaving or knitting process to produce the fire-protective woven fabrics (Fig. 21.1). In the case of fire-protective nonwoven fabrics, fibers are directly bonded (chemically, mechanically) together to produce the fabrics (Fig. 21.1). Sometimes, combustible fibers are directly used to manufacture woven or nonwoven fabrics, and these fabrics are chemically treated with FR substances to render them fire protective (Fig. 21.1). Fire-protective fabrics (as a single-layered woven fabric, or as a multilayered woven and/or nonwoven fabric system) are generally used as raw materials for

Thermal characterization of fire-protective fabrics

357

the manufacturing of workwear for firefighters, oil and gas industry workers, defense personnel, and/or cooks/chefs in order to provide them protection from a fire hazard while on duty (Bajaj and Sengupta, 1992; Rossi, 2003; Mandal et al., 2013; Shalev and Barker, 1983; Sun et al., 2000; Song et al., 2011a). Sometimes, single-layered fire-protective fabrics are also used in the sleepwear for children (in the age group of 10 months to 11 years) (Gordon and Pressley, 1978; Knudson et al., 1980). Therefore it is necessary to investigate the fire-protective performance of fabrics for providing the protection and safety to these wearers (Kahn et al., 2012; Song et al., 2016). Contextually, it is notable that the protective performance of fabrics depends upon the type of fire exposures faced by these wearers (Mandal et al., 2013). In general, these wearers frequently get exposed to radiant-heat, gas-based flame, and/or flash fire. Occasionally, first responders (e.g., firefighters, defense personnel) also get exposed to liquid-based flame (e.g., Molotov cocktail, petrol bomb) that is hurled by the protestors or rioters (Kemp et al., 2016). In addition, knees and elbows of crawled firefighters can come in contact with the hot solid surfaces and it causes significant burns on their bodies (Mandal and Song, 2014, 2016a,b). Fire-protective performance of fabrics can generally be measured in terms of their thermal stability and thermal insulation capacity (Song et al., 2016; Yip, 2014). Considering this, many standardized methods have been developed by different organizations [e.g., International Organizations for Standardizations (ISO), Comite´ Europe´en de Normalisation, American Society for Testing and Materials (ASTM), Canadian General Standards Board] for measuring the fire-protective performance of fabrics under different types of fire exposures such as radiant-heat, gas flame, combined radiant-heat and gas flame, and hot surface contact (Mandal et al., 2013; Rossi and Zimmerli, 1994). Some researchers also developed customized methods for measuring the fire-protective performance of fabrics under flash fire and liquid flame exposures. In the last few decades, many researchers used these standardized and/or customized methods for evaluating the fire-protective performance of fabrics. Based on these researches, it has been found that various factors related to fibers, yarns, fabrics, and/or fire exposures mainly affect the protective performance. In this chapter, various standardized and customized test methods for measuring the protective performance of fabrics under various fire exposures are thoroughly described. This section will help to understand and compare the differences among the test methods. Previous research on the fire-protective performance of fabrics is also systematically reviewed. This review will help to develop better understanding on the fire-protective performance. Various key issues related to the fire-protective performance of fabrics are also addressed. By solving these key issues, textile or materials researchers can develop new standards for measuring the fire-protective performance of fabrics. Also, they could engineer a new fire-protective fabric for high-performance workwear or sleepwear. These developments will provide a better occupational health and safety to the wearers such as firefighters, oil and gas industry workers, defense personnel, cooks/chefs, and/or children.

358

21.2

Thermal Analysis of Textiles and Fibers

Measurement of fire-protective performance of fabrics

In order to measure the fire-protective performance of fabrics, it is primarily required to determine the “thermal stability” and “thermal insulation capacity” of the fibers or fabrics (Bajaj, 1992; Song et al., 2016; Yip, 2014). The determination of the thermal stability will help to understand the integrity of the fabrics under fire exposures. Also, the determination of thermal insulation capacity will help to understand the adequacy of the fabrics for providing burn injury protection to the wearers for certain duration, especially when they are exposed to a fire hazard. For measuring the fire-protective performance, generally it is required to precondition the fibers or fabrics at 20 C 6 2 C and 65% 6 5% relative humidity for at least 24 hours as per ISO 139:2005. This preconditioning could help to obtain a fiber or fabric with defined moisture content (although this preconditioning is originally defined for natural fibers, it could be applicable for the synthetic fibers too).

21.2.1 Determination of thermal stability of fibers or fabrics For determining the thermal stability, it is necessary to understand the transitional characteristics [e.g., glass transitions (hard-to-soft transition), phase transitions (crystalline melting, melting), polymorphic transitions (changes in crystalline structure, thermal degradation), and property transitions (changes in heat capacity)] of the thermoplastic synthetic fibers or yarns used in a fabric under a fire exposure (Chen et al., 2006; Gaan et al., 2011a,b; Zhu et al., 2004). Thereafter, it is required to check the combustibility or flammability of the fabrics; actually, this will help to find out the ability of a fabric to ignite and propagate fire for burning or combustion (Bourbigot and Flambard, 2002; Yip, 2014).

21.2.1.1 Transitional characteristics test ISO 11357-1:2016 (2016) standard is used for understanding the transitional characteristics of any polymeric materials, including synthetic fibers/yarns. Especially to understand the transitional characteristics of a synthetic fiber or yarn, a dedicated ASTM D 7138:2016 (2016) standard is also available and widely used. As per these ISO and ASTM standards, differential scanning calorimetry (DSC) is used to detect the glass transition temperature, crystalline melting temperature, changes in heat capacity, and endothermic thermal degradation of the multiblended fiber or yarn specimens of 240 mg (Fig. 21.2). For a nonblended fiber or yarn a melting temperature device (MTD) can be used as per ASTM D 7138:2016 (2016) (Fig. 21.3). For this a specimen (2 mm length) of fiber or yarn is positioned in the MTD, and the temperature of the device is raised until it is not possible to visually detect the transition from a solid fiber state to a liquid fiber state. As the methods for evaluating the melting temperature of blended and nonblended fibers/yarns are different in ASTM D 7138:2016 (2016), results obtained from these two methods are

Thermal characterization of fire-protective fabrics

359

Specimen holder

Figure 21.2 ISO 11357-1:2016 (2016) and ASTM D 7138:2016 (2016) tester (DSC) at Empa, Switzerland.

Specimen holder Visual observation glass

Figure 21.3 ASTM D 7138:2016 (2016) melting temperature device at Empa, Switzerland.

incomparable. Also, the evaluation of melting temperature of nonblended fibers and yarns are dependent upon the visual observation of the experimenters; so, this MTD method is mainly suitable for the research and development purpose.

21.2.1.2 Flammability test ISO 4589-2:1996 (1996) standard can be used to check the flammability or combustibility of fabrics in terms of the minimum amount of oxygen required for its burning, that is, limiting oxygen index (LOI). For this, the upper end of a vertically oriented fabric specimen (80150 mm 3 10 mm) is ignited (using a gas flame igniter) in a mixture of oxygen and nitrogen flowing upward through a transparent chimney (Fig. 21.4). The minimum oxygen concentration required for the ignition is predicted by testing a series of specimens in different oxygen concentrations, and this minimum oxygen concentration is interpreted as the LOI of the fabric. Similar to ISO 4589-2:1996 (1996) standard, ASTM D 2863:2013 (2013) standard is

360

Thermal Analysis of Textiles and Fibers

Transparent chimney Specimen holder Gas flame igniter Oxygen and nitrogen supply point

Specimen

Figure 21.4 ISO 4589-2:1996 (1996) and ASTM D 2863:2013 (2013) tester at Empa, Switzerland.

Specimen holder Burner

Figure 21.5 ISO 15025:2016 (2016), ISO 6940:2004 (2004), and ISO 6941:2003 (2003) tester at Empa, Switzerland.

available to predict the LOI of a fabric specimen (80200 mm 3 1020 mm). In these ISO and ASTM standards the size of the specimens varies depending upon the thickness of the fabrics. ISO 15025:2016 (2016) standard is also used to check the flammability of fireprotective fabrics. As per this standard, a defined gas flame from a specified burner is exposed for 10 seconds to the surface or the bottom edge of a vertically oriented fabric specimen of 200 mm 3 160 mm (Fig. 21.5). From this exposure test, information on specimen’s fire propagation behavior [melting, spread of the flame, debris or flaming debris (material separating from the specimen during the test and falling from the specimen without or with flaming), hole (opening, break, or discontinuity with the original specimen), after flame time (duration of flaming after removal of the gas flame), afterglow (glowing or not after the removal of the gas flame), and/ or afterglow time (duration of afterglow)] is recorded. In order to understand the flammability of a multilayered fabric system, both sides of the system are exposed to the gas flame for 10 seconds. Also, the damaged length of the specimens is

Thermal characterization of fire-protective fabrics

361

recorded when the flame is exposed to the bottom edge of a vertically oriented fabric specimen. Similar to ISO 15025:2016 (2016), ISO 6940:2004 (2004), and @ISO 6941:2003 (2003) standards together can also be used to check the flammability of fire-protective fabrics; however, these two standards are mainly used for checking the flammability of nonfire-protective fabrics of different specimen sizes (200 mm 3 80 mm for ISO 6940:2004, 2004; 560 mm 3 170 mm for ISO 6941:2003, 2003). Similar to ISO 15025:2016 (2016) standard, ASTM D 1230:2016 (2016) standard is available to check the flammability of 45-degree aligned fabric specimen (50 mm 3 150 mm); however, this ASTM standard is not suitable to check the flammability of any fire-protective fabrics. As per the guideline of ISO 11612:2015 (2015), fabrics tested according to ISO 15025:2016 (2016) standard should meet the following criteria: fabric specimen should not permit the flame spread at any part of the lowest boundary of any flame to reach the upper or either vertical edge, give the flaming or molten debris, give the hole formation of 5 mm or greater, allow the afterglow and after flame time of $ 2 seconds.

21.2.2 Determination of thermal insulation capacity of fabrics In order to determine the thermal insulation capacity, fabrics are tested under various types of heat and fire exposures—radiant-heat, gas flame, flash fire, liquid flame, and hot surface contact (ISO 6942:2015, 2015; ISO 9151:2016, 2016; ISO 17492:2003, 2003; ISO 12127-1:2015, 2015; Kemp et al., 2016; Mandal et al., 2017). For these testing, one of the prime requirements is small and lightweight, inexpensive, and highly sensitive heat flux sensors that can properly simulate human (wearers) skin. To date, five different types of heat flux sensors (epoxy sensor, copper sensor, PyroCal copper sensor, aluminum sensor, skin simulant sensor) are available in the market and widely used for measuring the thermal insulation capacity of fabrics (Figs. 21.621.8). During the testing, this sensor is placed on the back side of the fabric (that will be aligned with the wearer’s skin), and the face side of the fabrics is exposed to fire. During and after the fire exposure, temperature increase of the sensor with respect to time is calculated by the thermocouple attached with the sensor and recorded in a computer. This temperature rise can further be used to calculate the heat flux through human skin (sensor) using different equations indicated by Mandal and Song (2015). From this heat flux, time to burns on human skin can be calculated based on Stoll curve or Henriques burn integral (HBI) equation (Mandal and Song, 2015). The temperature rise, heat flux, and/or time to burns can be used as the thermal insulation capacity of fabrics.

21.2.2.1 Radiant-heat insulation test ISO 6942:2015 (2015) standard is used for determining the thermal insulation capacity of a fire-protective fabric under the exposure of radiant-heat. According to this standard, a nonmoving vertically oriented fabric specimen is exposed to the radiant-heat of different levels of densities (low: 5 and 10 kW/m2; medium: 20 and 40 kW/m2; and high: 80 kW/m2). Here, radiant-heat is generated from six horizontally

Figure 21.6 Epoxy sensor manufactured by Precision Products, United States. Source: Courtesy Mr. Pat Kavanaugh (President, Precision Products, United States).

Figure 21.7 Copper sensors manufactured by Precision Products, United States. Source: Courtesy Mr. Pat Kavanaugh (President, Precision Products, United States).

Figure 21.8 PyroCal sensor manufactured by Precision Products, United States. Source: Courtesy Mr. Pat Kavanaugh (President, Precision Products, United States).

Thermal characterization of fire-protective fabrics

363

Radiant-heat source

Specimen support frame Heat shielding separator

Figure 21.9 ISO 6942:2015 (2015) and ASTM F 1939:2015 (2015) tester at Empa, Switzerland.

placed and electrically heated silicon carbide rods—these rods need a certain preheating time to reach a steady-state heat flux and that will be determined with a calibration measurement. A heat-shielding separator is placed in between the heated rods and specimen especially before and after the radiant-heat exposure to specimen (Fig. 21.9). As per the Method A of this standard, the specimen is exposed to the radiant-heat for a specific time, and following the exposure, the specimen is examined for visible changes. For method B, a copper sensor is placed on the technical back side of the specimen and the radiant-heat is exposed on the technical face side of the specimen. The temperature in the sensor to raise 12 C and 24 C is recorded in seconds. The means in seconds for three specimens is calculated as the heat transfer levels. The heat transmission factor (HTF) of the specimen is also reported in terms of the ratio of the transmitted to the incident heat flux densities. Similar to ISO 6942:2015 (2015), ASTM F 1939:2015 (2015) standard is available to determine the thermal insulation capacity of a fire-protective fabric under radiant-heat exposure. As per the guideline of ISO 11612:2015 (2015), the performance level of fabric specimens that are tested according to ISO 6942:2015 (2015) standard can be categorized as a: Level 1: 7 seconds # HTF , 20 seconds, Level 2: 20 seconds # HTF , 50 seconds, Level 3: 50 seconds # HTF , 95 seconds, and Level 4: HTF $ 95 seconds.

21.2.2.2 Gas flame insulation test ISO 9151:2016 (2016) standard is used for determining the insulation capacity of a fire-protective fabric under the exposure of gas flame. According to this standard, a nonmoving horizontally oriented fabric specimen (140 mm 3 140 mm) is exposed to 80 6 2 kW/m2 flame, that is, generated from a Meker propane gas burner placed beneath the specimen (Fig. 21.10). Before the flame exposure to the fabric, the Meker burner is calibrated at the heat flux of 80 6 2 kW/m2 by employing the nude sensor. Notably, a heat-shielding separator is placed in between the burner and specimen before and after the flame exposure. In order to measure the heat energy passing through the specimen during the flame exposure, a copper sensor is placed

364

Thermal Analysis of Textiles and Fibers

Heat-shielding separator Sensor

Burner flame Specimen holder

Figure 21.10 ISO 9151:2016 (2016) tester at Empa, Switzerland.

on top of and in contact with the specimen. The temperature in the sensor to raise 24 C 6 0.2 C is recorded in seconds. The mean seconds for three fabric specimens is calculated as heat transfer index (HTI), and this HTI is interpreted as the thermal insulation capacity of the fabric. Similar to ISO 9151:2016 (2016), ASTM D 4108:1987 (1987) standard was available for determining the thermal insulation capacity of fabrics under flame exposure. However, this standard is no more active for the testing purpose of fabrics. Recently, ISO 9151:2016 (2016) tester is also modified by replacing the copper sensor with a skin simulant sensor (Mandal, 2016). This newly modified tester can be effectively used for determining the thermal insulation capacity of different types of single- and multilayered fabrics. As per the guideline of ISO 11612:2015 (2015), the performance level of fabric specimens that are tested according to ISO 9151:2016 (2016) standard can be categorized as a: Level 1: 4 seconds # HTI , 10 seconds, Level 2: 10 seconds # HTI , 20 seconds, or Level 3: HTI $ 20 seconds.

21.2.2.3 Combined radiant-heat and gas flame insulation test ISO 17492:2003 (2003) is used for determining the thermal insulation capacity of a fire-protective fabric under the combined radiant-heat and gas flame exposure test. According to this standard, a nonmoving horizontally oriented fabric specimen (150 mm 3 150 mm) is exposed to combined radiant-heat (50%) and flame (50%) exposure of 80 6 2 kW/m2 (Fig. 21.11). The combined radiant-heat and flame is generated in the beneath of the specimen from 9 quartz T-500 infrared tubes and 2 Meker propane gas burners, respectively. Here, the gas burners are aligned at 45 degrees angle from the vertical and the quartz tubes are placed in the beneath and centered of the two burners. A heat-shielding separator is placed in between the burners/quartz tubes and specimen especially before and after the combined radiant-heat and flame exposure to the specimen. A copper sensor is placed on the top of the fabric specimen in order to measure the heat energy passing through the specimen during the exposure. The temperature in the sensor to raise 24 C 6 0.2 C

Thermal characterization of fire-protective fabrics

Specimen holder

Radiant-heat source

365

Meker burner

Heat-shielding separator

Figure 21.11 ISO 17492:2003 (2003), ASTM F 2703:2013 (2013), or ASTM F 2700:2013 (2013) tester at DuPont, Switzerland. Source: Courtesy Ms. Pauline Weisser (Application Development Specialist, DuPont, Switzerland).

is recorded in seconds. The mean seconds for three fabric specimens is calculated as HTI, and this HTI is interpreted as the thermal insulation capacity of the fabric. From this HTI, the time to generate second degree burn on human skin is also predicted from Stoll curve. Similar to ISO 17492:2003 (2003), ASTM F 2703:2013 (2013) (consider the stored energy inside the specimen after the fire exposure is ceased) and ASTM F 2700:2013 (2013) (did not consider the stored energy inside the specimen after the fire exposure is ceased) standards are also available for determining the thermal insulation capacity of fire-protective fabric under combined flame and radiant-heat exposures (ASTM F 2700, 2013; ASTM F 2703, 2013).

21.2.2.4 Flash fire insulation tests There are ISO 13506:2008 (2008) and ASTM F 1930:2015 (2015) standards for determining the thermal insulation capacity of a whole workwear under flash fire exposure. In these standards, workwear are put on a three-dimensional life-sized manikin, and this clothed manikin is exposed to the flash fire generated from several Meker gas burners. As these standards are for testing of whole workwear, these tests are expensive and cumbersome to carry out on regular basis. To date, no standardized methods are available for determining the thermal insulation capacity of fabrics used in the workwear under flash fire exposure. In order to determine the thermal insulation capacity of fabric under flash fire exposure, some standards (ISO 9151:2016, 2016; ISO 17492:2003, 2003; ASTM F 2703:2013, 2013; ASTM F 2700:2013, 2013) are widely used in Europe and North America. However, these standards can test only a small piece of horizontally aligned two-dimensional fabric specimen under flash fire exposure generated from one/two burner(s) and/or a radiant-heat source panel. As a result, these standards are far unrealistic in comparison to the ISO 13506:2008 (2008) or ASTM F 1930:2015 (2015) standard.

366

Thermal Analysis of Textiles and Fibers

Hexagonal panel

Sensor Gas burner

Figure 21.12 Flash fire insulation tester for calibration at Empa, Switzerland.

Fabric specimen (A)

(B)

Figure 21.13 Flash fire insulation tests for fabrics: (A) before flash fire exposure and (B) after flash fire exposure.

Considering the above-mentioned shortcomings, Empa-Switzerland has developed a customized method for determining the thermal insulation capacity of fabric under flash fire exposure (Mandal et al., 2017). According to this method, a hexagonal panel is instrumented with 36 epoxy sensors (developed by the Precision Products, United States) and placed in the middle of the 12 gas burners (Fig. 21.12). Flash fire is generated from these burners, and average flash fire intensity calculated from the 36 sensors is calibrated at 80 kW/m2 (Fig. 21.12). Then, six specimens (each specimen is 110 mm 3 1500 mm) of a fabric are clamped on the six sides of the hexagonal panel and engulfed into the 80 kW/m2 flash fire (Fig. 21.13). The duration of the flash fire is set at 4 and 8 seconds for the singleand multi-layered fabrics, respectively. During and 60 seconds (could be up to 120 seconds especially for multilayered fabrics) after the flash fire exposure, the temperature rise of the sensors is measured at every 0.1 seconds, and this temperature rise is further used for predicting the heat flux through the sensors as well as time required to generate the burns on the sensors.

Thermal characterization of fire-protective fabrics

367

Fuel resorvoir

Pulley system

Sensor Sensor board

Figure 21.14 Liquid flame insulation tester at Empa, Switzerland.

21.2.2.5 Liquid flame insulation test As no standardized method is available for determining the thermal insulation capacity of a fire-protective fabric under liquid flame exposure, Empa-Switzerland has developed a customized method for this (Kemp et al., 2016). As per this method, the upper edge of a fabric specimen (340 mm 3 170 mm) is clamped on a sensor board, and the lower edge is attached with a 500 g weight for providing a consistent tension on the specimen (Fig. 21.14). This sensor board have a size of 290 mm 3 180 mm and is made by a low thermal conductivity (0.18 W/m K), liquid (0.2% water regain at 23 C for 24 hours), and heat (up to 800 C)-resistant material; the sensor board is also instrumented with 10 epoxy sensors developed by the Precision Products, United States. For the test, a known amount of flammable liquid fuel is pipetted into a fuel reservoir that is positioned on the top of the specimen and then ignited. After the burning of the fuel for 10 seconds, it is tipped onto the specimen using a simple pulley system. The temperature rise of the 10 sensors during the exposure is recoded for 120 seconds, and from this temperature rise, heat flux and time to burn on human skin can be predicted using HBI software that is developed in-house.

21.2.2.6 Hot surface contact insulation test ISO 12127-1:2015 (2015) standard is used for determining the thermal insulation capacity of a fire-protective fabric in contact with hot surface. According to this standard, a heating cylinder and an aluminum sensor are mounted on a support frame. These heating cylinder and sensor are mounted in parallel faces, and their symmetrical axis is kept in a line (Fig. 21.15). A separator is positioned between the cylinder and the sensor for the heat-shielding. The temperature of the cylinder is controlled in between 100 C and 500 C, and a fabric specimen (of 80 mm diameter) is placed on the sensor. Then, the heat-shielding separator is removed, and the heating cylinder is lowered at a rate of 5 6 0.02 mm/s and placed in contact with the sensor. Notably, the contact force between the cylinder and sensor should be adjusted at 49 6 0.5 N by varying the weight on the cylinder. The temperature rise of the sensor is started to record when the distance between the cylinder and sensor

368

Thermal Analysis of Textiles and Fibers

Sensor

Heating cylinder Heat-shielding separator

Figure 21.15 ISO 12127-1:2015 (2015) tester at DuPont, Switzerland. Source: Courtesy Ms. Pauline Weisser (Application Development Specialist, DuPont, Switzerland).

is 10 mm, and the temperature recording is continued until the temperature of the sensor is 10 C higher than its starting value. The time between the start of timing and the moment when the temperature of the sensor is 10 C higher than its starting value is called as a threshold time (TT). Similar to ISO 12127-1:2015 (2015), ASTM F 1060:2016 (2016) standard is also available for determining the thermal insulation capacity of a fire-protective fabric under hot surface contact. As per the guideline of ISO 11612:2015 (2015), the performance level of fabric specimens that are tested according to ISO 12127-1:2015 (2015) standard can be categorized as a Level 1: 5 seconds # TT , 10 seconds, Level 2: 10 seconds # TT , 15 seconds, and Level 3: TT $ 15 seconds.

21.3

Evaluation of fire-protective performance of fabrics

Many researchers evaluated the fire-protective performance of fabrics in the last few decades. These researchers mainly evaluated the thermal stability and thermal insulation capacity of fibers and fabrics. These researches were carried out to understand the various factors that affect the thermal stability and thermal insulation capacity of fibers and fabrics. Notably, thermal stability is generally evaluated for single-layered fire-protective fabrics that are commonly used in sleepwear for children. However, evaluation of both thermal stability and thermal insulation capacity are required for single- and multilayered fabrics that are commonly used in workwear for firefighters, oil and gas industry labors, defense personnel, and/or cooks/chefs.

21.3.1 Thermal stability In order to evaluate the thermal stability, many researchers mainly investigated the transitional characteristics of noncombustible fibers: (1) chemically modified FR fibers (FR cotton, FR viscose, FR wool, FR polyester, FR nylon, FR acrylic),

Thermal characterization of fire-protective fabrics

369

and (2) inherently fire-resistant fibers [meta- and para-aramid, polyamide-imide, copolymer para-aramid, polybenzimidazole, polybenzazole, polypyridobisimidazole, others (phenolic, melamine, fluorocarbon, polyimide, and polyketone)] (Brown and Ennis, 1977; Bourbigot and Flambard, 2002; Nousiainen and MattilaNurmi, 1986; Powers and Serad, 1986; Varga et al., 2011; Wang et al., 2012a,b; Zhu et al., 2004; Zhu and Mao, 2014). Also, a group of researchers investigated the flammability of fire-protective fabrics (Baltusnikaite et al., 2006; Bourbigot and Flambard, 2002; Johnson, 1965; Ozcan et al., 2003; Ozcan et al., 2006; Salmeia et al., 2016). Contextually, Yip (2014) stated that the evaluation of both transitional characteristics and flammability is required for effectively evaluating the thermal stability of any fibers or fabrics.

21.3.1.1 Transitional characteristics Transitional characteristics of chemically modified FR fibers are investigated by many researchers using DSC/MTD (Nousiainen and Mattila-Nurmi, 1986; Varga et al., 2011; Wang et al., 2012a,b; Zhu et al., 2004; Zhu and Mao, 2014). It has been found that transitional characteristics of these fibers are mainly dependent upon the type and amount of FR substances applied on a substrate fiber [natural fibers (cotton, wool), regenerated fibers (viscose), or synthetic fibers (polyester, nylon, acrylic)]. Although halogen-based (e.g., fluorine, chlorine) FR substances can substantially improve the transitional characteristics of the substrate fiber, they are carcinogenic; eventually, their application is restricted especially in developed countries (Hirsch et al., 2017). Eco-friendly phosphorous-, nitrogen-, or siliconbased FR substance can improve the transitional characteristics of the substrate fiber (Salmeia et al., 2015, 2016; Serbezeanu et al., 2016); however, a combination of at least two of these substances can provide the better transitional characteristics results for any substrate fibers (Song et al., 2016). As indicated by Song et al. (2016), this better transitional characteristic is due to the synergistic effect of all the substances used. Some commercially available FR substances of this combined type (e.g., Proban and PyrovatexCP) may also drop the mechanical strength of a substrate fiber and that could affect the transitional characteristics of the fiber (Horrocks, 2011). Zhu et al. (2004) investigated the transitional characteristics of organophosphorus nitrogen-treated FR cotton fibers under air atmosphere. They mentioned that the FR substances promote dewatering and carbonizing reactions within the cotton fibers and that inhibit to produce the combustible L-glucose. This situation delays the transition of FR cotton fibers at low temperature. However, they indicated that the thermal degradation of the FR cotton fibers gradually occurs with increasing temperature (up to 400 C) and flammable aldehyde, ether, ketones, and/or a solid char residue are formed. Wang et al. (2012a,b) corroborated that the degradation of viscose fibers also behave similarly with increasing temperature and that the amount of solid residues increases with increasing the amount of phosphorous nitrogen based FR substances on the viscose fibers. Actually, similar chemical composition of cotton and viscose fibers resulted in this similar transitional characteristics

370

Thermal Analysis of Textiles and Fibers

CH2OH

CH2OH

H

OH H H OH

O

H

H OH

H

H

H OH

H OH H

OH

Figure 21.16 Chemical composition of cotton/viscose fibers. R2

H NH

CO C C

C NH

H

NH CO

H

CO

C R3

R1

Figure 21.17 Chemical composition of wool fiber (R 5 side chain of varying characteristics).

(Fig. 21.16). Contextually, Yip (2014) found that FR cotton and FR viscose fibers (manufactured by Lenzing AG, Austria) could lose B60% weight at 600 C in nitrogen atmosphere; however, these fibers could lose more than 80% of its weight under air atmosphere. Furthermore, Zhu and Mao (2014) stated that the degradation patterns of phosphorous, nitrogen, and chlorine-treated FR wool fabrics are significantly different from an untreated wool fabric. They found that phosphorous and chlorine could bond with the wool fibers and that results in insignificant mass losses and high char residues, especially after the exposure at 400 C. Contextually, both wool and cotton are originally natural fibers; however, the chemical composition of cotton and wool is very different (Figs. 21.16 and 21.17). As wool fiber contents high amount of nitrogen (16%), sulfur (3%4%), and water (standard moisture content 13.8%), its degradation temperature (570 C600 C) is very high in comparison to the degradation temperature of cotton fiber (300 C400 C). Eventually, transitional characteristics of a chemically modified FR wool fiber are inherently better than a cotton fiber. Although the natural/regenerated substrate fibersbased chemically modified FR fibers (FR cotton, FR viscose, FR wool) are widely used for fire-protective fabrics, some synthetic fibers (e.g., polyester, polyamide, acrylic) can also be used for this purpose. This is because these synthetic fibers have a low propensity for ignition or they slowly spared the fire if ignites. Recently, Varga et al. (2011) investigated the transitional characteristics of FR polyester fiber. It has been found that the degradation of this fiber started at 431 C, and this degradation process continuous up to 550 C. During this degradation, it was evident that this fiber could form a sticky substance and may adhere to the wearer’s skin. So, fire-protective fabrics consisting

Thermal characterization of fire-protective fabrics

371

of 100% of this fiber is not highly recommendable. If required, this fiber can be blended with other FR fibers (e.g., FR cotton, FR viscose); in this way, the dropping of sticky substances to the wearer’s skin can be eliminated by the char that is formed by the FR cotton or FR viscose (Nousiainen and Mattila-Nurmi, 1986). Similar to polyester, nylon and acrylic are also thermoplastic fibers; so, 100% of nylon and acrylic fibers may drop sticky substances on wearer’s skin. Due to this reason, the uses of synthetic substrate fiberbased chemically modified FR fibers are limited in the field of fire-protective fabrics. About four decades ago, Brown and Ennis (1977) thoroughly investigated the transitional characteristics of inherently fire-resistant meta- (commercially marketed as Nomex fiber developed by DuPont, United States) and para-aramid (commercially marketed as Kevlar fiber developed by DuPont, United States) fibers under air atmosphere. It was found that these fibers loses all its absorbed water and becomes dry at 100 C. And, at about 300 C, these fibers showed their glass transitions. Although meta-aramid fiber did not show any crystalline melting, it was confirmed that the thermal degradation of this fiber is generally started from 440 C. The crystalline melting and thermal degradation temperature of para-aramid fiber was 560 C and 590 C, respectively. Recently, Yip (2014) found that the meta-aramid (Nomex) fibers loses its 43%47% weight at 600 C under both the atmosphere of 100% nitrogen and 100% air. Similarly, paraaramid fibers (Kevlar, Twaron that is developed by Teijin Limited, Japan) loses it is B60% weight under nitrogen atmosphere; however, these fibers loses only B40% of its weight under air atmosphere. Like Nomex, commercially available meta-aramid fiber Conex (developed by Teijin Limited, Japan) shows similar transitional characteristics. Although chemical composition of Nomex (or Conex) and Kevlar (or Twaron) are similar, the molecular structure of these two fibers are very different (Fig. 21.18). As the amide groups in a Kevlar fiber are attached in the para positions of a benzene ring, the molecular chains of this fiber are parallel and symmetrical; these parallel and symmetrical molecular chains resulted in better transitional characteristics of Kevlar fiber in comparison to a Nomex fiber. Furthermore, a polyamide-imide fiber (commercially developed by Rhodia, France, and marketed as Kermel fiber) can also be classified as a meta-aramid fiber (Nomex or Conex); however, thermal conductivity of a commercially available Kermel fiber is much lower than the Nomex or Conex fiber at any O O C

CNH

NH O C

Nomex/Conex

O CNH

Kevlar/Twaron

Figure 21.18 Chemical structure of Nomex/Conex and Kevlar/Twaron fibers.

372

Thermal Analysis of Textiles and Fibers

H N

N

N

N

n

H

Figure 21.19 Chemical structure of PBI fiber.

temperature, and the heat capacity of these two fibers are significantly different. Also, a commercial copolymer para-aramid fiber such as Technora (developed by Teijin Limited, Japan) can be classified as an improved form of para-aramid fiber such as Kevlar; however, as this improvement is mainly related to its chemical stability, the transitional characteristics of Technora and Kevlar(or Twaron) is nearly same. Yip (2014) also indicated that the glass transition temperature (480 C) of the polybenzimidazole fiber (commercially developed by PBI Performance Products Incorporation, United States and marketed as PBI fiber) is significantly higher than the meta- and para-aramid fibers (Nomex, Conex, Kevlar, Twaron, Technora) under air atmosphere (Powers and Serad, 1986). Interestingly, polybenzimidazole fiber only loses 30% weight at 600 C temperature under nitrogen atmosphere. However, it loses nearly 70% of its weight under air atmosphere. In this context, Powers and Serad (1986) found that the degradation temperature of this fiber can be up to 1000 C under nitrogen atmosphere. As this fiber has a heterocyclic rigid rod-like molecular structure (Fig. 21.19), the degradation temperature of this fiber is generally higher than any aramid fibers. Based on this molecular structure, some other fibers are commercially developed [e.g., polybenzazole (commercially marketed as Zylon fiber that is developed by Toyobo Corporation, Japan), polypyridobisimidazole (commercially marketed as PPID or M5 fiber that is developed by Magellan Systems International, United States)] in the market; these fibers also possess similar transitional characteristics as polybenzimidazole. Along with the above-mentioned fibers (FR cotton, FR viscose, FR wool, FR polyester, meta- and para-aramid, polyamide-imide, copolymer para-aramid, polybenzimidazole, polybenzazole, polypyridobisimidazole), some other miscellaneous fibers can also be used for the manufacturing of fire-protective fabrics. These fibers are phenolic (commercially marketed as Kynol by Nippon Kynol Incorporation, Japan), melamine (commercially marketed as Basofil by BASF, Germany), fluorocarbon (commercially marketed as Teflon by DuPont, United States), polyimide (commercially marketed as P84 by Inspec Fibers, Austria), and polyketone (commercially marketed as PEEK by Zyex, United Kingdom) (Bajaj and Sengupta, 1992; Bourbigot and Flambard, 2002; Economy et al., 1973; Williams et al., 1985). The glass transition temperature of this miscellaneous fibers is generally within the range of 200 C300 C, and during the degradation of these fibers, these fibers can evolve different types of gasses (e.g., carbon dioxide, water vapor, nitride oxide, hydrogen cyanide), depending upon their chemical compositions.

Thermal characterization of fire-protective fabrics

373

21.3.1.2 Flammability Flammability of fire-protective fabrics in terms of their LOI values can vary in between 22% and 100%, and the inherently fire-resistant fibersbased fabrics generally possesses significantly high LOI in comparison to the chemically modified FR fibersbased fabrics (Bourbigot and Flambard, 2002; Song et al., 2016). Notably, LOI can only indicate the propensity of a fabric for ignition; and, a comparison between the LOI values of different fabrics indicate their time difference in delaying for ignitions. However, this indication may not give complete information related to the flammability of a fire-protective fabric (Johnson, 1965; Lee et al., 2014). For detail understanding of the flammability of any fire-protective fabrics, it is necessary to thoroughly investigate the fire propagation behavior the fabrics after ignition. Contextually, it has been found that the fire propagation behavior of a fabric is mainly dependent upon the physical characteristics of its constituent fiber(s) and warp/weft yarns (Baltusnikaite et al., 2006; Salmeia et al., 2016; Speece, 1974). Also, this fire propagation behavior is dependent upon the fabric parameters, namely, fabric type, count, design, and finish (Facts about Fabric Flammbility, 2003; Ozcan et al., 2003, 2006). The physical characteristics, size and form, of the constituent fiber(s) of a fabric have a significant effect immediately after its ignition (Johnson, 1965). As the size and form of synthetic fibers are more uniform in comparison to the natural fibers (Fig. 21.20), the surface of the synthetic fiberbased fabrics is generally smoother than the natural fibers based fabrics. Due to this surface smoothness, synthetic fibersbased fabrics did not offer a greater surface area and/or higher airsurface ratio. As the fire propagation behavior of a fabric is initially dependent on its surface combustion, synthetic fiberbased fabrics slowly propagate the fire. Also, the physical characteristics of the warp and weft yarns (twist, type, size, number of plies) used in a fabric play an important role in propagating the fire (Baltusnikaite et al., 2006). Generally, a highly twisted spun yarn traps less freely available air within its structure, and it helps to stop fire from propagating (Fig. 21.21). Depending upon the size and number of the plies of the yarns, the freely available air within a yarn may also differ, which could affect the flammability of the fabric. In this context, it is notable that yarns used for a fire-protective fabric should not be finished or coated with any combustible substance (lubricants

Figure 21.20 Size and form of (A) synthetic and (B) natural fibers (Arenas and Crocker, 2010).

374

Thermal Analysis of Textiles and Fibers

Twist in S direction High

Low

Figure 21.21 High and low twisted yarns. Fibers

Nonwoven

Yarn loop

Knitted

Warp

Weft

Woven

Figure 21.22 Different types of fabrics. Source: Courtesy Md. Sohanur Rahman Sobuj (Founder & Author at textilestudycenter.com).

or sizing ingredients). This kind of finish or coating could enhance the fire propagation behavior of the fabric (Ozcan et al., 2003). Although the physical characteristics of fibers and yarns are important, fabric parameters (type, count, design) play a crucial role in the context of propagating the fire (Baltusnikaite et al., 2006; Ozcan et al., 2003; Ozcan et al., 2003). Although a nonwoven fabric generally traps dead air within its constituent fibers, an availability of the free air inside this fabric could also help to propagate the fire (Fig. 21.22). If the size of the yarn loop within a knitted fabric or the interlacement between the warp and weft yarns in a woven fabric is highly open, it may allow the freely available air inside the fabric and that could enhance the fire propagation within the fabrics (Fig. 21.22) (Ozcan et al., 2003). Notably, the count (number of warp and weft yarns per inch) of a woven fabric could be too low; this low count fabric may increase the availability of the free air inside the fabric and propagate the fire. Variation in the designs of the woven fabrics may also vary the amount of openness and freely available air within the fabrics. This air can help to promote the propagation of the fire. For example, the openness of plain, twill, and statin weave fabrics are different. Eventually, the fire propagation behavior of these fabrics also varies (Fig. 21.23).

21.3.2 Thermal insulation capacity Thermal insulation capacity of a fabric significantly depends upon the types of exposed fires on the fabrics. Considering this, many researchers investigated the

Thermal characterization of fire-protective fabrics

Plain

2/2 Twill

375

5-Harness satin

Figure 21.23 Different designs of woven fabrics. Source: Courtesy Mr. Mazharul Islam Kiran (Editor-in-Chief and Founder of Textile Learner Blog, Bangladesh).

insulation capacity of fabrics under different types of fire exposures: radiant-heat, gas flame, combined radiant-heat and gas flame, flash fire, liquid flame, and hot surface contact (Lee and Barker, 1986; Mandal, 2016; Mandal et al., 2013; Mandal and Song, 2016a,b; Rossi and Zimmerli, 1994; Shalev and Barker, 1983). In these fire exposures, a fabric mainly insulates the heat energy transfer from the fire source to the wearers’ skin (sensors). Notably, physical properties of a fabric (weight, thickness, density, air permeability) affect its insulation capacity in a similar way under all types of fire exposures (Shalev and Barker, 1983; Mandal, 2016; Perkins, 1979; Sun et al., 2000). In general, a heavyweight, thick, and/or low densitybased fabric traps a lot of dead air within its structure, which can prohibit the heat energy transfer through the fabric. This situation can enhance the thermal insulation capacity of the fabric by increasing its thermal resistance. Also, a multilayered fabric system can trap a lot of dead air in between its constituent layers and that also increases the thermal insulation capacity of the fabric system. Contextually, the size of the microclimate region (air gap between the fabric and skin sensor) can substantially affect the thermal insulation capacity of the fabric especially under the exposure of radiant-heat, gas flame, or combined radiant-heat and gas flame (Song, 2007; Torvi, 1997; Wang et al., 2012a,b). An optimized size can trap a lot of dead air in the microclimate region and that can enhance the insulation capacity of the fabric. Notably, the moisture inside the fabric can change the thermal conductivity and heat capacity of the fabric, which can alter the thermal insulation capacity of the fabric (Rossi and Zimmerli, 1994; Lawson et al., 2004; Lee and Barker, 1986). As the heat transfer mechanism from the fire source to the wearers’ skin through fabrics is different under different types of exposures, the thermal insulation capacity of the fabrics also varies accordingly.

21.3.2.1 Insulation from radiant-heat Electromagnetic waves are mainly responsible for transferring the heat energy from a radiant-heat source to the fabric (Song et al., 2016). Here, the transfer of heat energy from the source to the fabric surface can be denoted by Eq. (21.1), where Q is the net heat energy transfer (kW/m2), σ is the StefanBoltzmann constant (5.670373 3 1028 W/m2/K4), A is the surface area of the fabric exposed (m2), εs is

376

Thermal Analysis of Textiles and Fibers

the effective emissivity of the radiant-heat source (dimensionless), εr is the emissivity of the radiant-heat source at temperature Tr ( C); that is, the ratio of source’s emissive power at Tr ( C) to that of black body at Tr ( C), αrfs is the absorptivity of fabric at Tr ( C) for incident waves from a black body at temperature Ts ( C); that is, the ratio of the absorptive power at Tr ( C) for incident waves from a black body at Ts ( C) to that of a black body at Tr ( C) for incident waves from a black body at Ts ( C). Q 5 σAεs ðεr Tr4 2 αrfs Ts4 Þ

(21.1)

When the heat energy transferred to the fabric surface, this energy transferred through the fabrics toward wearers’ skin. Therefore many researchers modeled the radiative-heat transfer through the fabrics by considering the absorptivity, transmissivity, and reflectivity of the fabrics (Torvi, 1997; Zhu et al., 2008). However, radiant emittance of the fabrics has not been considered in these models. During the radiant-heat exposure, temperature of the fabric increases. Due to this temperature raise, the radiant emittance of the fabric (radiant-heat emitted by a surface per unit area) increases as per the StefanBoltzmann Law [Eq. 21.2, where T is the temperature (K), j
is the radiant emittance (W/m2), ε is the emissivity of the fabric (dimensionless), and σ is the StefanBoltzmann Constant (5.670373 3 1028 W/m2/K4)]. In this state, fabric tends to absorb both long and short wavelength-based incident electromagnetic radiant-heat waves (i.e., generated from the radiant-heat source). This situation results in more emission of radiant-heat energy toward the sensor or wearer’s body [Eq. 21.3, where q is the radiant-heat energy (W) and A is the area of the fabric system (m2)], which ultimately lowers the thermal insulation capacity of the fabric. Notably, this emission is continued even after ceasing the exposure of radiant-heat on the fabrics. And, the radiant-heat emission can be occurred toward the wearers’ bodies (forward radiant emittance) and ambient environment (backward radiant emittance), depending upon the situation. Recently, Su et al. (2016) developed models for heat transfer through fabrics by considering its radiant emittance. The models applied during and after the radiant-heat exposures are shown in Eqs. (21.4) and (21.5), respectively. In these Eqs. (21.4) and (21.5), ρfab is the density of the fabric (g/cm3), (Cp)fab is the heat capacity of the fabric (J kg K), kfab is the thermal conductivity of the fabric [Eq. 21.6, where Vair% is the air percentage in the fabric’s pores; kfiber is the thermal conductivity of the air (W/m K), kfiber is the thermal conductivity of the fiber used in the fabric (W/m K)] at a temperature (T), t is the time (seconds), x is the distance (cm), qrad-trans1 is the transmitting portion of the incident radiant-heat on the fabric system, q1rad is the forward radiant emittance (W/m2), and q2rad is the backward radiant emittance (W/m2). 0


j 5 εσT 4

(21.2)

q 5 A 3 j
5 AεσT 4

(21.3)

Thermal characterization of fire-protective fabrics







  1  δT δ δT δqrad2trans1 δq rad δq2 rad 5 kfab ðTÞ 2 2 2 ρCp fab δt δx δx δx δx δx 

ρCp



  1  δT δ δT δq rad δq2 rad 5 k 2 2 ðTÞ fab fab δt δx δx δx δx

kfab ðTÞ 5 Vair %kair ðTÞ 1 ð1 2 Vair %Þkfiber ðTÞ

377

(21.4)

(21.5) (21.6)

Based on Eqs. (21.1)(21.6), it can be concluded that thermal insulation capacity of a fabric under radiant-heat exposure is dependent upon the surface area of the exposed fabrics, emissivity of the radiant-heat source and fabrics, density of the fabric, heat capacity of the fabric, thermal conductivity of the fabric, and radiant emittance of the fabric. Notably, surface emissivity of the fabric is significantly dependent on the surface optical properties of the fabric, namely, surface roughness and surface frictional coefficient. In addition, aluminized coating on the outer-layer (exposed to radiant-heat) of a multilayered fabric system (Frydrych et al., 2016; Hrynyk et al., 2013) and stored energy within the fabric system (Song et al., 2011a, b) significantly affect the thermal insulation capacity of the fabric.

21.3.2.2 Insulation from gas flame Under a gas flame exposure, hot gaseous molecules move toward the exposed fabric and form a boundary layer on the fabric surface (Song et al., 2016). At a particular point on the fabric surface, the thickness of the boundary layer (Δx) depends upon several factors namely v is the velocity of the gaseous molecules (m/s), ρ is the density of the gaseous molecules (g/cm3), and μ is the dynamic viscosity of the gaseous molecules (m2/s) (Eq. 21.7). ð Δx 5 ðv; ρ; μÞ

(21.7)

As the temperature of the boundary layer is higher than the temperature of the fabric surface, convective-heat energy transfer (Q) occurs from the boundary layer to the fabric surface. Here, convective-heat transfer coefficient (hc) is dependent on the several factors: v, ρ, μ, k is the thermal conductivity of the gaseous molecules (K m/W), cp is the specific heat capacity of the gaseous molecules (J/K), sg is the surface geometry of the fabric, and fc is the flow conditions of the gaseous molecules (Eq. 21.8). Depending upon the Reynolds number (Eq. 21.9), the flow conditions of the gaseous molecules can be laminar (Re , 5 3 105) or turbulent (Re . 5 3 105). If laminar, the heat energy (Q) transfer from the boundary layer to the fabric surface occurs axially (perpendicular to the fabric) and can be represented by Eq. (21.10), where Tα is the temperature of the hot gaseous molecules ( C), Ts is the temperature of the fabric surface ( C). If turbulent, the heat energy transfer

378

Thermal Analysis of Textiles and Fibers

from the boundary layer to the fabric surface both occurs axially and transversely; so, hc and P increases. ð hc 5 ðv; ρ; μ; k; cp ; fc ; sg ; fc Þ

(21.8)

ρvL μ

(21.9)

Q 5 hc ðTα 2 Ts Þ

(21.10)

Re 5

In this context, it is notable that gas flame comprises carbon particles of different diameters. During the flame exposure, these carbon particles also radiate some amount of heat onto the fabric surface. Therefore both convective- and radiativeheat impose on the fabric surface under the gas flame exposure. Considering this, Torvi (1997) investigated the heat energy transfer during the ASTM D 4108 (equivalent to ISO 9151) test with a single-layered fabric. He hypothesized that (1) convective- and radiative-heat energy transfer mainly occurs in the region 1, that is, between the burner and fabric, (2) conductive and radiative-heat energy transfer mainly occurs in the region 2, that is, within the fabric, and (3) convective- and radiative-heat transfer occurs in the region 3, that is, the air gap between the fabric and wearers’ skins. Based on this hypothesis, he modeled the radiativeheat transfer (qrad1 in kW/m2) together in the regions 1 and 3 as per Eq. (21.11), where σ is the StefanBoltzmann constant; εg, εf, and εb are emissivity of the hot gases, the fabric, and the burner head, respectively (dimensionless); Tg, Tf, Ta, and Tb are the temperatures of the hot gases, the fabric’s outer side, the ambient air, and the burner head, respectively ( C); Fa and Fb are view factors accounting for the geometry of the fabric with respect to the ambient air and the burner, and Af and Ab are the surface areas of the fabric and the burner head, respectively (m2). Also, the radiative-heat transfer (qrad2) in the region 2 can be modeled as per Eq. (21.12), where Ts, εs, and As are the firefighter’s body temperature ( C), emissivity (dimensionless), and surface area (m2), respectively, and Fs accounts for the geometry of the fabric with respect to the firefighter’s body. By using this concept, many researchers also further developed the model on heat energy transfer during the gas flame exposure using different conditions and fabrics namely after ceasing the flame exposure and multilayered fabrics (Ghazy and Bergstrom, 2010; Mell and Lawson 1999; Sawcyn and Torvi, 2009; Torvi and Threlfall, 2006; Zhu and Zhou, 2013). qrad1 5 σεg Tg4 2 σεf Fa ð1 2 εg ÞðTf4 2 Ta4 Þ 1

σFb ð1 2 εg ÞðTb4 2 Tf4 Þ


 

 1 1 Fb ð1 2 εg Þ 1 2 εf =εf 1 Af =Ab ð1 2 εb Þ=εb

(21.11)

Thermal characterization of fire-protective fabrics

379

σðTf4 2 Ts4 Þ 




  qrad2 5
ð1 2 εs Þ=εs 1 As =Af 1=Fs 1 1 2 εf =εf

(21.12)

Based on Eqs. (21.7)(21.12), it can be concluded that thermal insulation capacity of a fabric under gas flame exposure is dependent upon the features of the gaseous molecules: velocity, density, dynamic viscosity, temperature, thermal conductivity, specific heat capacity, flow conditions (laminar or turbulent). Along with the features of the gaseous molecules, characteristics of the fabrics (surface geometry, surface temperature) also play an important role on the thermal insulation capacity. Contextually, it has also been found that heat shrinkage and after wash effect of the fabric (Atalay et al., 2015; Crown et al., 2002; Rossi and Zimmerli, 1994) substantially lowers its thermal insulation capacity. The thermal insulation capacity of a fabric can be improved by assembling a superabsorbent fabric in the multilayered fabric system (Houshyar et al., 2015).

21.3.2.3 Insulation from combined radiant-heat and gas flame Many researchers investigated the thermal insulation capacity of fabrics under the combined exposure of radiant-heat and gas flame (He and Li, 2015; Lee and Barker, 1986; Li et al., 2012; Shalev and Barker, 1984). In this combined exposure, radiant- and convective-heat transfer occurs from the fire source to the fabric surface as per the mechanism described in the Section 21.3.2.1 and 21.3.2.2. Eventually, the emissivity of the radiant-heat source, features of the gaseous molecules in the flame (velocity, density, dynamic viscosity, temperature, thermal conductivity, specific heat capacity, flow conditions), and characteristics of the fabric (exposed surface area, emissivity, surface geometry, surface temperature) substantially affect the thermal insulation capacity of fabric under this combined exposure. Considering this, Udayraj et al. (2017) recently modeled the thermal insulation capacity (ts) of the fabrics (in terms of their time to generate second degree burn on wearers’ bodies) as per Eq. (21.13), where kf is the thermal conductivity of fiber (W/m K), ρf is the density of the fiber (g/cm3), cp,f is the specific heat of the fiber, Lf is the fabric thickness (cm), Ad is the fabric aerial density (g/cm3), ρa is the density of the air within fabric (g/cm3), ka is the conductivity of the air within fabric (W/m K), cp,a is the specific heat of the air (J/K), Lg is the width of the air gap (cm), Lf is the fabric thickness (cm), and qv is the incident heat flux from the combined radiant-heat and flame source (kW/m2). kf ts Ad ρa ka cp;a Lg ρ2 f c3 p;f L3 f q00 5f ; ; ; ; ; 2 ρf cp;f L f ρf Lf ρf kf cp;f Lf k3 f

! (21.13)

Contextually, Shalev and Barker (1984) found that the thermal insulation capacity of the fabric is nearly same under the combined exposure of radiant-heat and flame with a ratio of 50:50 and 70:30. However, they found that the heat energy

380

Thermal Analysis of Textiles and Fibers

transfer through the fabrics is quite different under the radiant-heat and flame exposure. The radiant-heat mainly penetrate through the pores exists in the fabrics; however, convective-heat is mainly impose on the fabric surface and then blown through it. Nevertheless, they concluded that the heat transfer through fabrics under this combined exposure is a complex phenomenon of absorption, reradiation, conduction, and perhaps forced convection. Lee and Barker (1986) also indicated that thermal insulation capacity of fabric under this combined exposure is always higher than 100% radiant-heat exposure. Wang et al. (2012a,b) and Li et al. (2012) also indicated that the amount of moisture in the microclimate region also significantly affect the thermal simulation capacity of the fabrics. Recently, He and Li (2015) found that a multilayered fabric system can store heat energy during the exposure to the combined radiant-heat and flame. As a result, the transmission of the heat energy toward the wearers’ bodies gets significantly reduced, and the thermal insulation capacity of the fabric becomes high. However, this stored energy dissipated toward the wearers’ body when the exposure is ceased or during the cooling of the fabric and that lowers the thermal insulation capacity of the fabric. He and Li (2015) mentioned that the stored energy did not affect the thermal insulation capacity immediately when the exposure is ceased. After 940 seconds of cooling, the stored energy dissipated toward the wearers bodies and lowers the thermal insulation capacity of the fabric. If the exposure time is long, the stored energy inside the fabric system is high, and this energy starts to dissipate toward the wearers’ body immediately when the exposure is ceased and lowers the thermal insulation capacity.

21.3.2.4 Insulation from flash fire It has been observed that the flash fire is generally a combination of radiant-heat, flame, fire balls, and hot gases (ASTM F 1930:2015, 2015; ISO 13506, 2008, 2008; Song, 2004). As the flash fire comprises radiant-heat and flame, the radiative- and convective-heat transfer mechanism from the fire source to the fabric surface could similar as described in the preceding sections. In addition, fire balls could significantly damage or degrade the surface of the fabrics, which will ultimately affect the radiative- and convective-heat transfer through the fabrics. Notably, hot gases could blow through the pores exists on the fabric surface and move toward the wearers’ body. This situation will reduce the thermal insulation capacity of the fabric. Recently, Mandal et al. (2017) studied the thermal insulation capacity of fabrics under flash fire exposure. They found that thermal insulation capacity of the fabric is generally higher under the flash fire exposure in comparison to the individual flame or radiant-heat exposure of 80 kW/m2. It has been observed that flash fire is mainly a combination of flame, radiant-heat, hot gasses, and fire balls. Due to this combined effect, the heat energy did not directly penetrate through the fabrics; rather, the heat energy mainly deflected from the surface of the fabrics. This deflection of thermal energy enhances the insulation capacity of the fabric. Notably, the configurations of the source of the flash fire (angle, distance from the fabric) have significant effect on the thermal insulation capacity of the fabric. Notably, it has been found that stored

Thermal characterization of fire-protective fabrics

381

energy inside the fabrics is very high under the flash fire exposure, and this stored energy significantly affects the thermal insulation capacity of the fabric. Mandal et al. (2017) extensively studied the thermal insulation capacity of the fabrics under flash fire exposure; however, they did not study the effect of moisture and microclimate region on the thermal insulation capacity of fabrics. Also, it may require to thoroughly investigating the stored energy inside the fabrics and modeling the heat transfer through fabrics. Research on these existing limitations may develop a comprehensive knowledge on the thermal insulation capacity of fabrics under flash fire exposure.

21.3.2.5 Insulation from liquid flame Under the liquid flame exposure, the liquid substance could transmit through the fabric depending upon the repellency and permeability/penetrability of the fabrics (Kemp et al., 2016). If the repellency or permeability/penetrability of a fabric is high, it will transmit more liquids through its structure. Eventually, it will help to expose the greater amount fabric surface to the flame and that will reduce the thermal insulation capacity of the fabric. Recently, Kemp et al. (2016) found that the flow of the liquid could vary depending upon the fiber content of a fabric. If a fabric contents more natural fiber such as cotton, it absorbs more liquid and possesses a low thermal insulation capacity. In addition, if the surface of the fabrics is finished or coated by some chemical substances, the liquid may disperses/splits from the fabric surface. This kind of dispersion/splitting may enhance the thermal insulation capacity of the fabric. Depending upon this dispersion rate, the fabric surface may also get damaged and that can affect the thermal insulation capacity of the fabric. Although the characteristics of the fabric are important, the volume of the liquid also plays a significant role for the thermal insulation capacity. If the volume of the liquid is very high, a greater amount of liquid may disperses from the fabric surface which increases the thermal insulation capacity. Also, a low volume of liquid may not cause enough flame on the fabric surface and increases the thermal insulation capacity of the fabric. The angle of incidence of the liquid is also crucial for the thermal insulation capacity. If the angle is very high, there is a chance to split the liquid from the fabric surface and enhances the thermal insulation capacity. Based on the preceding discussion, it can be confirmed that the features of the liquids and characteristics of the fabrics can substantially affect the thermal insulation capacity of the fabrics. Notably, some other factors (moisture content of the fabric, microclimate between the fabric and sensor) could also affect the thermal insulation capacity; however, no research has been carried out to date on these factors.

21.3.2.6 Insulation from hot surface contact During the exposure to the hot surface contact, conductive-heat energy [q0 (t)] transfer occurs from the hot surface to the one side of a fabric surface (surface 1 that is

382

Thermal Analysis of Textiles and Fibers

exposed to hot surface) as per Eq. (21.14) (Mandal and Song, 2016a,b). Then, this heat energy gradually transmitted [qv(t)] toward the opposite side of the fabric (surface 2 that is in contact with the wearers’ skin) as per Eq. (21.15). From this side of the fabric, heat energy gradually transmits [q0 v(t)] toward the wearers’ skin as per Eq. (21.16). In Eqs. (21.14)(21.16), ΔPF1 is the temperature difference between the hot surface plate and fabric surface 1 ( C), ΔXp is the thickness of the hot surface plate (cm), kP is the thermal conductivity of the hot surface plate (W/m K), A is the area of contact between hot surface plate and fabric surface 1 (m2), hPF1 is the thermal conductance coefficient between the hot surface plate and fabric surface 1 (W/m/K), ΔXF is the thickness of the fabric (cm), VA is the fabric’s air volume (cm3), VF is the fabric’s total volume (cm3), kγ is the thermal conductivity of fabric’s solid phase (W/m K), kα is the thermal conductivity of fabric’s solid phase (W/m K), P is the fabric porosity (dimensionless), TF is the temperature of the fabric ( C), E is the rate of heat energy generated per unit volume of fabric (J/m3), ργ is the density of the solid fiber phase of the fabric (g/cm3), ρα is the density of the gaseous air phases of the fabric (g/cm3), CPγ is the specific heat values of the solid fiber phase of the fabric (J/K), CPα is the specific heat of the gaseous air phase of the fabric (J/K), A0 is the area of contact between the fabric and skin (cm2), Xs is the thickness of the sensor (cm), and ks is the thermal conductivity of the sensor (W/m K). q0 ðtÞ 5

ΔPF1 


   ΔXP =ðkP AÞ 1 1=ðhPF1 AÞ 1 ΔXF = 1 2 VA =VF kγ 1 VA =VF kα A (21.14)

q00 ðtÞ 5 2 kγ ð1 2 PÞA 2 ρα CPα PA

@TF @TF @TF 2 kα PA 1 EA@XF 2 ργ CPγ ð1 2 PÞA dXF @XF @XF @XF

@TF dXF @XF

(21.15) q000 ðtÞ 5

ΔXF =





ΔF2S 
   1 2 VA =VF kγ 1 VA =VF kα A0 1 1=ðhF2S A0 Þ 1 ΔXS =ðkS A0 Þ (21.16)

From Eqs. (21.14)(21.16), it can be concluded that several factors affect the heat energy transfer through fabrics toward the skin during the hot surface contact exposure. Eventually, these factors have considerable effect on the thermal insulation capacity of the fabrics. Although these factors are well established, it is possible that some other factors may affect the heat energy transfer through fabrics in the presence of moisture. However, no research has been carried out to understand the thermal insulation capacity of the moistened fabrics under the exposure to hot surface contact.

Thermal characterization of fire-protective fabrics

21.4

383

Key issues related to the fire-protective performance of fabrics

ISO 11357-1:2016 (2016) and ASTM D 7138:2016 (2016) are available and used for evaluating the transitional characteristics of different types of fibers. However, a comprehensive study to compare the differences in transitional characteristics of different types of fibers at a typical, standardized, and controlled temperature (that is generally exposed by workers firefighters, oil and gas industry workers, defense personnel, and cooks/chefs) has not been carried out to date. This comparison would help to understand the best suitable fibers for use in a particular type of workwear. ISO 15025:2016 (2016), ISO 6940:2004 (2004), and ISO 6941:2003 (2003) standards are available and used to check the fire propagation behavior of fabrics. However, a comprehensive study on grading the fire propagation behavior of different fire-protective fabrics is still scanty. Considering this, a set of commonly used fire-protective fabrics should be tested in future using ISO 15025:2016 (2016) standard, and the images of the fire exposed surface of these fabrics can be collected. By utilizing the image analysis software, these pictures could be graded as per Likert scale. In future, this scale can be used for grading the flammability of any type of fire-protective fabrics. A great deal of standards (ISO 6942:2015, 2015; ASTM F 1939:2015, 2015; ISO 9151:2016, 2016; ASTM F 2700:2013, 2013; ASTM F 2703:2013, 2013; ISO 121271:2015; ASTM F 1060:2016, 2016) is available to determine the thermal insulation capacity of fire-protective fabrics under radiant-heat, gas flame, and/or hot surface contact. Also, there are customized methods for determining the thermal insulation capacity of fabric under flash fire and liquid flame exposure. However, these standards can simulate the fire exposures with an intensity of up to 80 kW/m2 or 500 C. As firefighters can be exposed to high-intensity fire ( . 100 kW/m2) especially when they work in a structural fire, there is a need to develop the standards that can simulate very high-intensity fire exposures. Moreover, firefighters and oil gas industry worker can be exposed to the low-intensity fire exposures (e.g., flame, radiant-heat) for long duration. However, no standards are available for determining the thermal insulation capacity of fabrics under low-intensity long duration fire exposure. An availability of these test standards could help to understand the change in fabric properties with the exposures. As the moisture and stored energy inside the fabrics under fire exposures are prime factors to affect the thermal insulation capacity (Lawson et al., 2004; Song et al., 2011b), some new standard methods are needed to develop (or modify the existing methods) by considering these factors. Also, presently available standard methods for determining the thermal insulation capacity of fabrics are cumbersome and difficult to carry out on a routine basis. To date, no empirical model has been developed for predicting the thermal insulation capacity of fabrics under different types of fire exposures: radiant-heat, gas flame, flash fire, liquid flame, and hot surface contact. Recently, Mandal and Song (2014) attempted to develop the empirical multiple linear regression and artificial neural network models for predicting the thermal insulation capacity of fabric under

384

Thermal Analysis of Textiles and Fibers

radiant-heat, gas flame, and hot surface contact exposures. However, these models were developed by using a small data set. In future, it is necessary to validate and extend these models for new and large data sets. Finally, fabrics for workwear available in the market are mainly heavyweight and thick. Although this type of fabrics is good for providing protection to the wearers under fire exposures, it can exert a lot of heat stress on on-duty wearers’ bodies. Eventually, it causes a great discomfort to the wearers, reduces working performance, and even induces heat stress. In future, it is necessary to develop the lightweight fabric for the workwear by using the latest technologies: for example, smart technology, nanotechnology. This type of newly developed fabrics could lower the discomfort to the wearers.

21.5

Summary and conclusion

Fire-protective fabrics are used as raw materials for the manufacturing of workwear (for firefighters, oil and gas industry workers, defense personnel, and cooks/chefs) and children sleepwear. For providing the effective protection and safety to these wearers, it is necessary to measure the fire-protective performance of fabrics. In order to measure the fire-protective performance of fabrics, it is required to determine the thermal stability and thermal insulation capacity of the fabrics. In order to determine the thermal stability and thermal insulation capacity of the fabric, many standardized methods are available on the market: ISO 11357-1:2016 (2016), ASTM D 7138:2016 (2016), ISO 15025:2016 (2016), ISO 6940:2004 (2004), ISO 6941:2003 (2003), ISO 6942:2015 (2015), ASTM F 1939:2015 (2015), ISO 9151:2016 (2016), ASTM F 2700:2013 (2013), ASTM F 2703:2013 (2013), ISO 12127-1:2015, and ASTM F 1060:2016 (2016). Furthermore, researchers from Empa-Switzerland have developed two customized methods for determining the thermal insulation capacity of the fabrics under the flash fire and liquid flame exposures. In the last few decades, many researchers used these standardized and customized methods for evaluating the fire-protective performance. Based on their studies, it has been found that many factors related to the fibers, yarns, fabrics, and fire exposures could affect the fire-protective performance of fabrics. Although a lot of studies have been carried out for evaluating the fire-protective performance of fabrics, there are still some unattended and not fully understood key issues related to the fire-protective fabrics. In future, these key issues need to be resolved in order to develop new standards or fabrics for workwear and children sleepwear. These developments could provide the better protection and safety to the workers and children across the world.

References Arenas, J.P., Crocker, M.J., 2010. Recent trends in porous sound absorbing materials. Sound Vib. 44, 12. ASTM D 4108:1987, 1987. Standard Test Method for Thermal Protective Performance of Materials for Clothing by Open-Flame Method.

Thermal characterization of fire-protective fabrics

385

ASTM D 2863:2013, 2013. Standard Test Method for Measuring the Minimum Oxygen Concentration to Support Candle-Like Combustion of Plastics (Oxygen Index). ASTM F 2703:2013, 2013. Standard Test Method for Unsteady-State Heat Transfer Evaluation of Flame Resistant Materials for Clothing With Burn Injury Prediction. ASTM F 2700:2013, 2013. Standard Test Method for Unsteady-State Heat Transfer Evaluation of Flame Resistant Materials for Clothing With Continuous Heating. ASTM F 1939:2015, 2015. Standard Test Method for Radiant Heat Resistance of Flame Resistant Clothing Materials With Continuous Heating. ASTM F 1930:2015, 2015. Standard Test Method for Evaluation of Flame Resistant Clothing for Protection Against Fire Simulations Using an Instrumented Manikin. ASTM D 7138:2016, 2016. Standard Test Method to Determine Melting Temperature of Synthetic Fibers. ASTM D 1230:2016, 2016. Standard Test Method for Flammability of Apparel Textiles. ASTM F 1060:2016, 2016. Standard Test Method for Evaluation of Conductive and Compressive Heat Resistance (CCHR). Atalay, O., Bahadir, S.K., Kalaoglu, F., 2015. An analysis on the moisture and thermal protective performance of firefighter clothing based on different layer combinations and effect of washing on heat protection and vapour transfer performance. Adv. Mater. Sci. Eng. Available from: https://doi.org/10.1155/2015/540394. Bajaj, P., 1992. Bull. Mater. Sci. 15 (1), 67. Bajaj, P., Sengupta, A.K., 1992. Textile Progress 22 (24), 1. Baltusnikaite, J., Suminskeine, R., Milasius, R., 2006. Mater. Sci. 12 (2), 167. Bourbigot, S., Flambard, X., 2002. Fire Mater. 26, 155. Brown, J.R., Ennis, B.C., 1977. Text. Res. J. 47 (1), 62. Chen, S., Zheng, Q., Ye, G., Zheng, G., 2006. J. Appl. Polym. Sci. 102, 698. Crown, E.M., Dale, J.D., Bitner, E., 2002. Fire Mater. 26 (45), 207. Economy, J., Wohrer, L.C., Frechette, F.J., Lei, G.Y., 1973. Appl. Polym. Symp. 21, 8189. Facts About Fabric Flammbility, 2003. ,http://comfycozy.com/fed_flammability.pdf.. Fabric for Ideas, n.d. ,http://www.trevira.com/fileadmin/download/broschueren_cs_gesamt/ gesamt_gb.pdf.. Frydrych, I., Cichocka, A., Gilewicz, P., Dominiak, J., 2016. Polymers 8 (10), 348. Gaan, S., Mauclaire, L., Rupper, P., Salimova, V., Tran, T., Heuberger, M., 2011a. J. Anal. Appl. Pyrolysis 90 (1), 33. Gaan, S., Salimova, V., Rupper, P., Ritter, A., Schmid, H., 2011b. Flame retardant functional textiles. In: Pan, N., Sun, G. (Eds.), Functional Textiles for Improved Performance. Woodhead Publishing, United Kingdom, pp. 98130. Ghazy, A., Bergstrom, D.J., 2010. Numer. Heat Transfer A—Appl. 58 (9), 702. Gordon, P.G., Pressley, T.A., 1978. Burns 5 (1), 13. Hirsch, et al., 2017. Arch. Toxicol. 91 (1), 407. Horrocks, A.R., 2011. Polym. Degrad. Stab. 96 (3), 377. Houshyar, S., Padhye, R., Troynikov, O., Nayak, R., Ranjan, S., 2015. J. Text. I 106 (12), 1394. Hrynyk, R., Frydrych, I., Irzmanska, E., Stefko, A., 2013. Text. Res. J. 83 (17), 1860. He, J., Li, J., 2015. Text. Res. J. 86 (3), 235244. ISO 4589-2:1996, 1996. Plastics  Determination of Burning Behaviour by Oxygen Index. ISO 6941:2003, 2003. Textile Fabrics  Burning Behavior  Measurement of Flame Spread Properties of Vertically Oriented Specimens. ISO 17492:2003, 2003. Clothing for Protection Against Heat and Flame  Determination of Heat Transmission on Exposure to Both Flame and Radiant Heat. ISO 6940:2004, 2004. Textile Fabrics  Burning Behavior  Determination of Ease of Ignition of Vertically Oriented Specimens.

386

Thermal Analysis of Textiles and Fibers

ISO 13506:2008, 2008. Protective Clothing Against Heat and Flame  Test Method for Complete Garments  Prediction of Burn Injury Using an Instrumented Manikin. ISO 11612:2015, 2015. Protective Clothing  Clothing to Protect Against Heat and Flame  Minimum Performance Requirements. ISO 6942:2015, 2015. Protective Clothing  Protection Against Heat and Fire  Method of Test: Evaluation of Materials and Material Assemblies When Exposed to a Source of Radiant Heat. ISO 12127-1:2015, 2015. Clothing for Protection Against Heat and Flame  Determination of Contact Heat Transmission Through Protective Clothing or Constituent Materials. ISO 11357-1:2016, 2016. Plastics  Differential Scanning Calorimetry (DSC). ISO 15025:2016, 2016. Protective Clothing  Protection Against Flame  Methods of Test for Limited Flame Spread. ISO 9151:2016, 2016. Protective Clothing Against Heat and Flame  Determination of Heat Transmission on Exposure to Flame. Johnson, A.J., 1965. A Comparison of the Flammability of Three Types of Fabrics Used in Girls’ Sleepwear (M.Sc. thesis). Oklahoma State University, Stillwater, OK. Kahn, S.A., Patel, J.H., Lentz, C.W., Bell, D.E., 2012. J. Burn. Care. Res. 33 (1), 152. Kemp, S.E., Annaheim, S., Rossi, R.M., Camenzind, M.A., 2016. Test method for characterising the thermal protective performance of fabrics exposed to flammable liquid fires. Fire Mater. Available from: https://doi.org/10.1002/fam.2416. Kilinc, F.S., 2013. Handbook of Fire Resistant Textiles. Elsevier Science and Technology, United Kingdom. Knudson, M.S., Bolieu, S.L., Larson, D.L., Brown, R.D., 1980. Burns 6 (4), 255260. Lawson, L.K., Crown, E.M., Ackerman, M.Y., Dale, J.D., 2004. Int. J. Occup. Saf. Ergon. 10 (3), 227238. Lee, Y.M., Barker, R.L., 1986. J. Fire Sci. 4 (5), 315. Lee, C.H., Salit, M.S., Hassan, M.R., 2014. Adv. Mater. Sci. Eng. Available from: https:// doi.org/10.1155/2014/514036. Li, X., Lu, Y., Li, J., Wang, Y., Zhou, L., 2012. Fiber Polym. 13 (4), 549. Liang, S., Neisius, N.M., Gaan, S., 2013. Prog. Org. Coat. 76 (11), 1642. Mandal, S., 2016. Studies of the Thermal Protective Performance of Textile Fabrics Used in Firefighters’ Clothing Under Various Thermal Exposures (Ph.D. thesis). University of Alberta, Edmonton, Canada. Mandal, S., Song, G., 2014. Ann. Occup. Hyg. 58 (8), 1065. Mandal, S., Song, G., 2015. Text. Res. J. 85 (1), 101. Mandal, S., Song, G., 2016a. AATCC J. Res. 3 (2), 8. Mandal, S., Song, G., 2016b. J. Ind. Text. Available from: https://doi.org/10.1177/ 0123456789123456. Mandal, S., Song, G., Ackerman, M., Paskaluk, S., Gholamreza, F., 2013. Text. Res. J. 83 (10), 1005. Mandal, S., Annaheim, S., Pitts, T., Camenzind, M., Rossi, R., 2017. Text. Res. J. Available from: https://doi.org/10.1177/0040517517723020. Mell, W.E., Lawson, J.R., 1999. Fire Technol. 36 (1), 39. Nousiainen, P., Mattila-Nurmi, M., 1986. J. Appl. Polym. Sci. 31 (2), 597. Neisius, M., Stelzig, T., Liang, S., Gaan, S., 2014. Flame retardant finishes for textiles. In: Paul, R. (Ed.), Functional Finishes for Textiles. Elsevier, United Kingdom, pp. 429461. Ozcan, G., Dayioglu, H., Candan, C., 2003. Text. Res. J. 73 (10), 883. Ozcan, G., Dayioglu, H., Candan, C., 2006. Indian J. Fibre. Text. Res. 31, 330. Perkins, R.M., 1979. Text. Res. J. 49 (4), 202.

Thermal characterization of fire-protective fabrics

387

Powers, E.J., Serad, G.A., 1986. History and development of polybenzimidazoles. In: Seymour, R.B., Kirshenbaum, G.S. (Eds.), High Performance Polymers: Their Origin and Development. Springer, The Netherlands, pp. 355373. Rossi, R., 2003. Ergonomics. 46 (10), 1017. Rossi, R., Zimmerli, T., 1994, January. Influence of humidity on the radiant, convective and contact heat transmission through protective clothing materials. In: Fifth International Symposium on Performance of Protective Clothing: Improvement through Innovation, San Francisco, CA. Salmeia, K.A., Fage, J., Liang, S., Gaan, S., 2015. Polymers 7 (3), 504. Salmeia, K.A., Gaan, S., Malucelli, G., 2016. Polymers 8 (9), 1. Salmeia, K.A., Jovic, M., Ragaisiene, A., Rukuiziene, Z., Milasius, R., Mikucioniene, D., et al., 2016. Polymers 8 (8), 293. Sawcyn, C.M.J., Torvi, D.A., 2009. Text. Res. J. 79 (7), 632. Serbezeanu, D., Butnaru, I., Varganici, C., Bruma, M., Fortunato, G., Gaan, S., 2016. RSC Adv. 6, 38371. Shalev, I., Barker, R.L., 1983. Text. Res. J. 53 (8), 475. Shalev, I., Barker, R.L., 1984. Text. Res. J. 54 (10), 648654. Song, G., 2004. Modeling Thermal Protection Outfit for Fire Exposures (Ph.D. thesis). North Carolina State University, Raleigh, NC. Song, G., 2007. J. Ind. Text. 36 (3), 193. Song, G., Paskaluk, S., Sati, R., Crown, E.M., Dale, J.D., Ackerman, M., 2011a. Text. Res. J. 81 (3), 311. Song, G., Cao, W., Gholamreza, F., 2011b. Text. Res. J. 81 (11), 120. Song, G., Mandal, S., Rossi, R., 2016. Thermal Protective Clothing: A Critical Review. Elsevier Science and Technology, United Kingdom. Speece, J., 1974. EC74-492 Fabric Flammability and Clothing. Historical Materials from University of Nebraska-Lincoln Extension. Su, Y., He, J., Li, J., 2016. J. Ind. Text. Available from: https://doi.org/10.1177/ 1528083716644289. Sun, G., Yoo, H.S., Zhang, X.S., Pan, N., 2000. Text. Res. J. 70 (7), 567. Torvi, D.A., 1997. Heat Transfer in Thin Fibrous Materials Under High Heat Flux Conditions (Ph.D. thesis). University of Alberta, Edmonton, Canada. Torvi, D.A., Threlfall, T.G., 2006. Fire Technol. 42 (1), 27. Udayraj, Talukdar, P., Das, A., Alagirusamy, R., 2017. J. Text. I 108 (2), 260. Varga, K., Noisternig, M.F., Griesser, U.J., Aljaz, L., Koch, T., 2011. Lenzinger Berichte 89, 50. Wang, X., Li, Q., Di, Y., Xing, G., 2012a. Fiber Polym. 13 (6), 718. Wang, Y.Y., Lu, Y.H., Li, J., Pan, J.H., 2012b. Fiber Polym. 13 (5), 647. Williams, L.L., Updegrapff, I.H., Petropulos, J.C., 1985. Amino resins. In: Tees, R.D., Pohelein, G.W. (Eds.), ACS Symposium Series. American Chemical Society, pp. 11011115. Yip, P.W., 2014. Analysis of Two Methods for Characterization of Flame Resistant Military Fabrics and Commercial Textile Fibers: Simultaneous DSC-TGA and Pyrolysis GCMS. U.S. Army Natick Soldier Research, Development and Engineering Center. Zhu, J., Mao, N., 2014, May. Keratin polymer having improved thermal properties. In: 14th World Textile Conference AUTEX 2014, Bursa, Turkey. Zhu, F.L., Zhou, Y., 2013. Fibres Text. East. Eur. 21 (5), 85. Zhu, P., Sui, S., Wang, B., Sun, K., Sun, G., 2004. J. Anal. Appl. Pyrolysis 71, 645655. Zhu, F., Zhang, W., Song, G., 2008. Fire Saf. J. 43 (6), 401409.

Acronyms and abbreviations

Abbreviations 3HB 3HH AFM AN ATRP BCF CF Ch Cr DCP DEA DETA DMA DMAc DMF DMS DMTA DSC DTA F-DSC FT-IR ɣ-BL K-CL K-VL GDD GL HEMA I IA iPP LC LOI LVDT MA MDSC MMT MTMA MWCNT

3-hydroxybutyric acid 3-hydroxyhexanoic acid atomic force microscopy acrylonitrile atom transfer radical polymerization bulk continuous fiber carbon fiber cholesteric crystal, crystalline di-cumyl peroxide dielectric analysis dielectric thermal analysis (the trademark of Polymer Labs) dynamic mechanical analysis dimethyl acetamide dimethyl formamide dimethyl sulfoxide dynamic mechanical thermal analysis (the trademark of Polymer Labs) differential scanning calorimetry, differential scanning calorimeter differential thermal analysis fast-scan DSC Fourier-transform infrared spectroscopy ɣ-butyrolactone K-caprolactone K-valerolactone 1,3-diglycerol diacrylate glycoside 2-hydroxyethyl methacrylate isotropic itaconic acid isotactic polypropylene liquid crystal, liquid crystalline limiting oxygen index linear voltage differential transformer methyl acrylate modulated DSC (the trademark of TA Instruments) montmorillonite modulated TMA (the trademark of TA Instruments) multiwalled carbon nanotube

390

MT-DSC MT-TMA N NMR Nylon M5T OM PAEK PAN PBI PBS PBT PCL PCR PE PEG PEEK PEKK PEN PET PGA PHA PHO PLA PLLA PDLA POSS PP PPT PTFE PU PVDF Q-DSC S, Sm SAXS SEM TA TEM TG TGA TMA TSA UHMMPE UHMWPE VAc WAXS XRD

Acronyms and abbreviations

modulated temperature DSC modulated temperature TMA nematic nuclear magnetic resonance poly(2-methylpentamethylene terephthalamide) optical microscopy poly(aryl ether ketone) polyacrylonitrile poly(benzimidazole), poly[2,20 -(m-phenylene)-5,50 -bis-benzimidazole] poly(butylene succinate) poly(butylene terephthalate) polycaprolactone polymerase chain reaction polyethylene poly(ethylene glycol) poly(ether ether ketone) poly(ether ketone ketone) poly(ethylene naphthalate), poly(ethylene-2,6-naphthalene dicarboxylate) poly(ethylene terephthalate) poly(glycolic acid) polyhydroxyalkonoate poly(3-hydroxyoctanoate) poly(lactic acid) poly(L-lactic acid) poly(D-lactic acid) polyhedral oligomeric silsesquioxanes polypropylene poly(propylene terephthalate) poly(tetrafluoroethylene) polyurethane poly(vinylidene fluoride) quasi-isothermal DSC smectic small angle X-ray scattering scanning electron microscopy thermal analysis transmission electron microscopy thermogravimetry thermogravimetric analysis thermomechanical analysis thermal stress analysis ultrahigh molecular mass polyethylene ultrahigh molecular weight polyethylene vinyl acetate wide-angle X-ray scattering X-ray diffraction

Acronyms and abbreviations

391

Acronyms β CR CLTE CVTE Cp Cv D D ΔCp ΔCp100% ΔHc ΔHf ΔHf DDR DR ‘ ΔSf DR dSh/dT dTex HR HFlA HRA HTF HTI HTL ΔT E’ EO E\ E” fAM fX k λ Lf Mn p Q Q q1rad ρ RAF T

isothermal compressibility cooling rate coefficient of linear thermal expansion coefficient of volumetric thermal expansion heat capacity at constant pressure heat capacity at constant volume denier, the unit of linear mass density of fibers (g per 9000 m of the fiber) used mostly in the United States and the United Kingdom (1 D 5 9 Tex, 0.9 dTex) dye diffusion coefficient change of specific heat capacity at the glass transition change of specific heat capacity at the glass transition for a 100% amorphous sample heat of crystallization heat of fusion equilibrium heat of fusion (heat of fusion at the equilibrium melting point) drawdown ratio draw ratio length entropy change on melting draw ratio rate of shrinkage the unit of linear density of fibers, 1 dTex 5 10 Tex heating rate heat flow rate amplitude heating rate amplitude heat transmission factor heat transfer index heat transfer level temperature difference dynamic storage tensile modulus dynamic storage tensile modulus parallel to the draw direction dynamic storage tensile modulus perpendicular to the draw direction dynamic loss tensile modulus orientation function of the chains in the amorphous phase orientation function of the chains in the crystalline phase thermal conductivity draw ratio fabric thickness number average molecular mass pressure heat convective-heat energy transfer forward radiant-emittance density rigid amorphous fraction temperature

392

Tα tan δ Tg Tex σ Tblock Tco Tcp Tcco Tccp Td Tm  Tm Tmp Tp Tr, TR Ts, TS τi U V Vair X

Acronyms and abbreviations

temperature of α-relaxation loss factor (5Ev/E0 ) glass transition temperature the unit of linear mass density of fibers (g per 10 km of the fiber) used mostly in Canada and continental Europe (1 D 5 9 Tex, 0.9 dTex) stress block temperature starting temperature of crystallization peak temperature of crystallization starting temperature of cold crystallization peak temperature of cold crystallization degradation temperature melting point equilibrium melting point peak temperature of melting peak temperature reference temperature sample temperature induction time of crystallization internal energy volume air percentage in fabric’s pores crystallinity

Index

Note: Page numbers followed by “f” and “t” refer to figures and tables, respectively. A Acrylic fibers coefficient of linear thermal expansion (CLTE) of, 85 86, 86f mechanical properties of, 299t production of, 298 thermal analysis of, 297 thermal shrinkage, thermomechanical analysis of, 88f Acrylics, 205 Additive effects on polyglycolides, 158, 159f on polyhydroxyalkanoates, 174 175, 179f, 181f, 182f on polylactic acid, 164 166, 165f on poly-ε-caprolactone, 172 174, 173f Aliphatic nylon fibers, thermal analysis of, 223 Brill transition, 227 cold crystallization, 226 227, 227f decomposition, 230 drawing, 229t glass transition, 226, 226t melting point, 225 226, 226t, 228 229, 229t reorganizational behavior, 227 228, 228f semicrystallinity, 229 transcrystallinity, 230 Aliphatic polyesters, thermal properties of, 151 poly-ε-caprolactone, 166 174 additives, effects of, 172 174, 173f copolymers, 170 172, 172t homopolymers, 166 170, 167f, 168f polyglycolides, thermal properties of, 153 158 additives, effects of, 158, 159f copolymers, 156 158, 157f

homopolymers, 155f, 158 polyhydroxyalkanoates, 174 182 additives, effects of, 174 175, 179f, 181f, 182f copolymers, 175 176, 177f homopolymers, 174 175, 174f polylactic acid, 158 166 additives, effects of, 164 166, 165f copolymers, 162 164, 163f homopolymers, 158 161, 160f, 162f p-Aramid, chemistry of, 328f As-spun PET fibers DMA analysis of, 100 101, 101f first heating cooling second heating of, 53 56, 55f thermal shrinkage, thermomechanical analysis of, 88, 89f, 90f Avrami equation, 119 120 B BHET. See Bis(hydroxyethyl)terephthalate (BHET) Bis(hydroxyethyl)terephthalate (BHET), 135 136 Blending method, 114 115, 114f Bulking, 14 C Carbon fibers mechanical properties of, 299t production of, 298 stabilization process, 298 299 thermal analysis of, 297 Celluloses, 108 109 Cellulosic yarns crystallinity, TGA analysis of, 73, 74f CLTE. See Coefficient of linear thermal expansion (CLTE)

394

Coefficient of linear thermal expansion (CLTE), 81, 84 85, 87 of acrylic yarns, 85 86, 86f of PET yarns, 85 86 Collagens, 109 Copolymers poly-ε-caprolactone, 170 172, 172t poly(glycolic acid-co-L-lactic acid), 260 261 polyglycolides, 156 158, 157f polyhydroxyalkanoates, 175 176, 177f polylactic acid, 162 164, 163f Cotton chemical treatment effect on thermal stability, TGA analysis of, 78, 78f heat capacity of, 119f CP spin-draw polymer, 136 Crystallinity differential scanning calorimetry of, 53 60, 55f, 57f, 58f, 59f, 60f of natural fibers, DSC measurement of, 111 115 blending method, 114 115, 114f glass transition, 112 114, 113f of natural polymers, thermal analysis characterization of, 119 122, 120f of poly(vinyl alcohol), 277 279, 280f, 281t TGA analysis of, 73, 74f D Die swell, 8 Differential scanning calorimetry (DSC), 17 basics of, 17 23 definition of, 17 of fibers, 36 67 crystallization, 53 60, 55f, 57f, 58f, 59f, 60f flame-retardant systems, effect of, 62 67, 62f, 63f, 64f, 65f, 66f glass transition of, 49 53, 49f, 50f, 51f, 52f, 53f on heating, 23 25 melting behavior, 41 45, 41f, 42f, 43f, 44f, 46f, 47 49, 47f, 48f Nylon M5T, 37, 38f shrinkage effect, 39 41, 40f thermal stability, 60 62, 61f

Index

specific heat capacity calibration (heat flow calibration), 31 36 heat-flux, 17 19, 18f of liquid crystalline polymers, 332 333, 332f modulated temperature, 20 natural fibers, crystallinity of, 111 115 blending method, 114 115, 114f glass transition, 112 114, 113f parameters, 20 23, 22f of polyacrylonitrile fibers, 303 312, 303f, 304f, 304t, 306f, 307t of polybenzimidazole fiber, 293, 294f of polyetheretherketone, 250, 253, 254f of polyetherketoneketone, 254 255, 255f, 256f of poly(vinyl alcohol) fibers, 279, 287 289, 288f power-compensation, 19 20, 19f temperature calibration of, 23 36 on cooling, 26 30, 26f, 27f, 28f, 29f energy (enthalpy) calibration, 30 Differential thermal analysis (DTA), 17 19, 18f Dilatometry, 81 82 Dimethylacetamide, 293 DMA. See Dynamic mechanical analysis (DMA) Drawing, 2, 3f, 13 14 effect on melt spun PVA fibers, 279, 282f hot, of poly(vinyl alcohol) fibers, 274 275, 281t of polypropylene fibers, 209f pseudohexagonal, 205 206, 209 212 Dry spinning of poly(vinyl alcohol) fibers, 273 DTA. See Differential thermal analysis (DTA) Dynamic mechanical analysis (DMA), 81, 95 of aliphatic nylon fibers, 228f, 230 autotension, 96 97 of liquid crystalline polymers, 330 331, 332f melt spinning of, 9 10 of nylon-6, 235t, 236 237, 236f, 237t of nylon-6,6, 239 240 of polyacrylonitrile fibers, 311 312, 313f of polybenzimidazole fiber, 293

Index

of polyethylene fibers, 197 198 pretension, 96, 97f storage modulus, 98 100, 99f stress strain curve in tensile mode, 97, 98f tensile storage modulus, 100f, 101 102, 101f, 102f, 103f Dyneema UHMW highly drawn fiber, modulated temperature TMA curves of, 197 198, 198f E Elastin, 109 Electrospinning of nylon-6, 237 239, 238f of nylon-6,6, 241 F Fast-scanning chip calorimetric technique, 116 117 FC. See Ferromagnetic core (FC) Ferromagnetic core (FC), 83 84 Fiber process structure property relationships, 7 drawing, 13 14 heat setting and bulking, 14 spinning, 8 13, 9f, 10f, 11f, 12f Fire-protective fabrics evaluation of, 368 382 flammability, 373 374, 373f, 374f, 375f thermal insulation capacity, 374 382 thermal stability, 368 374, 370f, 371f, 372f key issues related to performance, 383 384 manufacturing process of, 356f thermal characterization of, 355 thermal insulation capacity, 361 368, 362f combined radiant-heat and gas flame insulation test, 364 365, 365f flash fire insulation tests, 365 366, 366f gas flame insulation test, 363 364, 364f hot surface contact insulation test, 367 368, 368f liquid flame insulation test, 367, 367f

395

radiant-heat insulation test, 361 363, 363f thermal stability, 358 361 flammability test, 359 361, 360f transitional characteristics test, 358 359, 359f Flame-retardant systems effect on fibers differential scanning calorimetry of, 62 67 Flame retardation effect on thermal stability, TGA analysis of, 78, 79f Flammability of fire-protective fabrics, 359 361, 360f evaluation of, 373 374, 373f, 374f, 375f Flash fire insulation of fire-protective fabrics, 365 366, 366f, 380 381 Fourier-transform infrared spectroscopy (FTIR) oxidative stabilization of PAN fiber by, 317 321, 318f, 318t, 319f, 320f Free-to-shrink, 39, 40f G Gas flame insulation of fire-protective fabrics, 363 364, 364f, 377 379 Gel-drawing, 95 Gel-spinning, 95 of poly(vinyl alcohol) fibers, 273 274 Gel-spun PE fibers differential scanning calorimetry of, 199 200 thermal properties of, 200 from UHMWPE, 199 Glass transition temperature differential scanning calorimetry of, 49 53, 49f, 50f, 51f, 52f, 53f of natural fibers, 112 114, 113f of polybenzimidazole fiber, 293 of polypropylene fibers, 215, 217f, 218f of poly(vinyl alcohol), 275 H Heat-flux DSC, 17 19, 18f Heat setting, 14 Heat transfer index (HTI), 363 364 Homopolymers poly-ε-caprolactone, 166 170, 167f, 168f polyglycolides, 155f, 158 polyhydroxyalkanoates, 174 175, 174f

396

Homopolymers (Continued) polylactic acid, 158 161, 160f, 162f Hot surface contact insulation of fireprotective fabrics, 367 368, 368f, 381 382 HTI. See Heat transfer index (HTI) Hydrated inorganic salt, 338, 339f K Keratins, 108 Kevlar/Twaron fiber, chemical composition of, 371f L Lag-shear model, 328 329 Latent heat, 336 343, 345 347, 349 350 LCPs. See Liquid crystalline polymers (LCPs) Lignin, 109 Limiting oxygen index (LOI), 359 360 Linear voltage or variable differential transformer (LVDT), 83 84, 84f Liquid crystalline polymers (LCPs) chemical structure of, 326 328 functionality, 327f PET versus, 327f process structure property relationships of, 328 329, 329t tensile modulus, 325 326, 328 330 thermal analysis of, 325, 330f, 331f, 332f Liquid flame insulation of fire-protective fabrics, 367, 367f, 381 LOI. See Limiting oxygen index (LOI) M Melting behavior differential scanning calorimetry of, 41 45, 41f, 42f, 43f, 44f, 46f, 47 49, 47f, 48f of natural fibers, 115 117, 116f of polypropylene fibers, 211 215, 212f, 213f, 214f, 216f Melt spinning, 8 9, 9f, 12 13 morphology development, 10 11, 10f of polypropylene fibers, 207, 208f of poly(vinyl alcohol) fibers, 273, 277, 279f spun birefringence, 12 13, 12f structure formation, 11 12, 11f

Index

Mercury dilatometer, 81 82 Modulated temperature DSC (MT-DSC), 20 heat flow calibration, 36 Modulated-temperature TMA (MT-TMA), 81, 84 85 of an as-spun and drawn PET single filament, 85f Morphology of poly(vinyl alcohol), 277, 278f MT-DSC. See Modulated temperature DSC (MT-DSC) N Natural fibers characterization, TGA analysis of, 77 78, 77f structure of, 106 109 silk proteins, 107 108 keratins, 108 celluloses, 108 109 collagens, 109 lignin, 109 elastin, 109 thermal analysis of. See Natural fibers, thermal analysis of Natural fibers, thermal analysis of, 105 crystallinity, 111 115 blending method, 114 115, 114f glass transition, 112 114, 113f crystallization kinetics, 119 122, 120f melting behavior, 115 117, 116f natural polymer bound water systems, 122 124, 123f natural polymer metal ions systems, 124 126, 125f specific heat capacity, 117 119 experimental measurement, 119, 119f theoretical prediction, 117 119, 118f Natural polymer bound water systems, thermal analysis study of, 122 124, 123f Natural polymer metal ions systems, thermal analysis study of, 124 126, 125f Nomex/Conex fiber, chemical composition of, 371f Nylon-6 crystallinity, 234 DSC melting curves of, 43 45, 46f

Index

dynamic mechanical analysis, 235t, 236 237, 236f, 237t electrospinning, 237 239, 238f polymorphs, physical properties of, 233t reorganizational behavior, 233 234, 234t thermal behavior of, 232 239 zone-annealing method, 235 236 zone-drawing method, 235 236 Nylon-6,6 dynamic mechanical analysis of, 239, 240 electrospinning, 241 thermal behavior of, 239 241, 240f thermal shrinkage of, 240 241, 241f Nylon 66 fiber crystallization of, 58f glass transition of, 52 53, 53f melting behavior of, 58f Nylon fibers definition of, 224 hydrogen bonding in, 225, 225f M5T fibers, DMA analysis of, 101 102, 102f nomenclature of, 224 225 producers of, 224f production and basics, 223 225 O Organic hydrocarbons, 338 341, 339f, 340f Oxidative stabilization (cyclization), of PAN fiber, 300, 312 321 chemical composition, effect of, 313 315, 314f by Fourier-transform infrared spectroscopy (FT-IR), 317 321, 318f, 318t, 319f, 320f processing parameters, effect of, 315 317 P PAEK. See Poly(aryl ether ketone) (PAEK) fibers PBI. See Polybenzimidazole (PBI) fiber PCR. See Polymerase chain reaction (PCR) PE. See Polyethylene (PE) fibers PEEK. See Poly(ether ether ketone) (PEEK) PEG. See Polyethylene glycol (PEG) PEK. See Poly(ether ketone) (PEK) PEKK. See Poly(ether ketone ketone) (PEKK)

397

PEN. See Poly(ethylene naphthalate) [poly (ethylene-2,6-naphthalene dicarboxylate), PEN] PGA. See Polyglycolic acid (PGA) Phase-change material (PCM), 335 336 capsulated, structure of, 338f thermal management performance of, 342 347, 344f, 345f, 346f, 347f types of, 337 342 hydrated inorganic salt, 338, 339f organic hydrocarbons, 338 341, 339f, 340f polyethylene glycol, 341 342, 341f working principle of, 336 337, 337f m-Phenylene, 293, 294f Physiological comfort, 335 PLGA. See Poly(lactic-co-glycolide) (PLGA) PLLA. See Poly(L-lactide) acid (PLLA) Polyacrylonitrile, thermal analysis of, 301 302 Polyacrylonitrile (PAN) fiber, 297 carbonization of, 299 300 oxidative stabilization (cyclization) of, 300, 312 321 chemical composition, effect of, 313 315, 314f by Fourier-transform infrared spectroscopy (FT-IR), 317 321, 318f, 318t, 319f, 320f processing parameters, effect of, 315 317 oriented fibers, structure of, 300 301 production of, 301 radiation effect on pyrolysis, 75 77, 76f thermal analysis of, 303 312 DMA characterization, 311 312, 313f DSC characterization, 303 312, 303f, 304f, 304t, 306f, 307t heating rates in air and nitrogen, TGA measurements of, 306 309, 307f, 309f, 309t PAN/IA/MA, PAN/IA, and PAN/VAc fibers, TGA measurements of, 309 311, 310t, 312f thermal stabilization of, 300 Poly(aryl ether ketone) (PAEK) fibers, 247 commercial data, 249t producers of, 248 249

398

Poly(aryl ether ketone) (PAEK) fibers (Continued) resins, structure of, 248t synthesis of, 248 thermal analysis of, 247 Polybenzimidazole (PBI) fiber, 291 β-relaxation, 293 chemical structure of, 372f differential scanning calorimetry of, 293, 294f DMA of, 293 glass transition temperature, 293 short-term thermo-oxidative properties of, 292t Poly(butylene terephthalate), thermal analysis of, 147 Poly-ε-caprolactone (PCL), thermal properties of, 166 174 additives, effects of, 172 174, 173f copolymers, 170 172, 172t homopolymers, 166 170, 167f, 168f Polydioxanone sutures (PPDX), 265 degradation of, 267f, 268f hydrolytic degradation of, 265 266 morphological changes in, 265 Poly(ether ether ketone) (PEEK) crystallization of, 62f, 63f differential scanning calorimetry of, 250, 253, 254f Hoffman Weeks plots, 250 251, 251f melt spinning of, 253 254 structure of, 248t synthesis of, 248 thermal analysis of, 249 254 thermal stability of, 62, 63f thermogravimetric analysis, 251, 252f Poly(ether ketone) (PEK) structure of, 248t thermal analysis of, 250 251 Poly(ether ketone ketone) (PEKK) differential scanning calorimetry of, 254 255, 255f, 256f structure of, 248t thermal analysis of, 254 257 thermogravimetric analysis, 255 257, 256f, 257f Polyethylene (PE) fibers, 197 drawability of, 201, 202f dynamic mechanical analysis of, 197 198

Index

fixed length, 200 201 free-to-shrink, 200 201 macromolecules, 198 199 melting behavior of, 203 pseudohexagonal, 200, 203 thermal properties of, 200 Polyethylene glycol (PEG), 341 342, 341f semicrystallinity of, 261 Poly(ethylene naphthalate) [poly(ethylene-2,6-naphthalene dicarboxylate), PEN], 191 DMA analysis of, 100, 100f fixed length, 193 194 free-to-shrink, 193 195, 195f, 196f melt spinning of, 192 193 polycondensation of, 191 192 thermal analysis of, 147, 193 Poly(ethylene terephthalate) (PET) chip, characterization of, 136 137 coefficient of linear thermal expansion (CLTE) of, 85 86 crystallization of, 50f, 55f DMA analysis of, 98 100, 99f fiber spinning of, 8 flame retardation effect on thermal stability, TGA analysis of, 78, 79f glass transition of, 51f history of, 135 melting behavior of, 42, 44f physical characteristics of, 138, 138t polymerization of, 135 136 processing of, 137 138 single filament, MT-TMA measurement of, 85f thermal shrinkage, thermomechanical analysis of, 87, 91f, 92f thermal stability of, 61f Poly(glycolic acid-co-L-lactic acid) copolymers, 260 261 Poly(lactic-co-glycolide) (PLGA) hydrolytic degradation of, 266 Poly(L-lactide) acid (PLLA) biodegradation of, 263 264, 263f crystallinity of, 264 265 degradation of, 264 265 drawing mechanism of, 262 dry spinning of, 264 265 glass-transition temperature of, 264 melting point of, 261, 262f, 264

Index

semicrystallinity of, 261 Polyester fibers, 133 thermal analysis of, 139 145, 139f, 140f, 141f, 142f, 143f poly(butylene terephthalate), 147 polyethylene naphthalate, 147 polypropylene terephthalate, 147 in the 21st century, 145 147, 146f Polyglycolic acid (PGA) chemical structure of, 153f thermal properties of, 153 158 additives, effects of, 158, 159f copolymers, 156 158, 157f homopolymers, 155f, 158 Polyhydroxyalkanoates (PHAs), thermal properties of, 174 182 additives, effects of, 174 175, 179f, 181f, 182f copolymers, 175 176, 177f homopolymers, 174 175, 174f Polylactic acid (PLA), thermal properties of, 158 166 additives, effects of, 164 166, 165f copolymers, 162 164, 163f homopolymers, 158 161, 160f, 162f Polymerase chain reaction (PCR), 117 Polymer flammability, TGA analysis of, 73 75, 75f, 76f Polymerization of PET, 135 136 Polypropylene (PP) fibers, 205 crystallization of, 57f, 59f, 205 206 drawing of, 209f endothermic hysteresis peaks of, 49 52, 52f geotextiles, 207 glass transition of, 52 53, 53f hollow fibers, 207 manufacturing of, 207 mechanical properties of, 207 melting behavior of, 45, 47f, 48f, 59f melt spinning of, 207, 208f morphology of, 10 nonwovens, 206 207 processing of, 207 208 properties of, 208 211 resistance to chemicals, 206 semicrystalline, 205 206 thermal analysis of, 211 219 glass transition, 215, 217f, 218f

399

melting, 211 215, 212f, 213f, 214f, 216f results/experimental conditions, comparison of, 216 219, 219f thermal conductivity of, 206 Polypropylene terephthalate, thermal analysis of, 147 Polypropylene yarns, surface-to-volume ratio, TGA analysis of, 72 73, 72f Polystyrene, glass transition of, 49f Poly(vinyl alcohol) (PVA) fibers manufacturing of, 272 275 dry spinning, 273 gel spinning, 273 274 hot drawing and stabilization, 274 275 melt spinning, 273 wet spinning, 272 273 thermal analysis of, 271 crystallinity, 277 279, 280f, 281t differential scanning calorimetry of, 279, 287 289, 288f dynamic mechanical analysis, 282 287, 283f, 284f, 285f glass transition temperature, 275 melt spinning, 277, 279f morphology, 277, 278f residual moisture effect, 275 277, 277f thermogravimetric analysis, 275, 276f thermomechanical analysis, 279, 282f wet spinning, 277 zone drawing, 286f, 287 289 Power-compensation DSC (PC-DSC), 19 20, 19f PP. See Polypropylene (PP) fibers PPDX. See Polydioxanone sutures (PPDX) Process Structure Property Performance Tetrahedron, 2f PVA fibers, DMA analysis of, 100 101, 101f R Radiant-heat insulation of fire-protective fabrics, 361 363, 363f, 375 377 Radiation effect on pyrolysis, TGA analysis of, 75 77, 76f RAF. See Rigid amorphous fraction (RAF) Residual moisture effect, on poly(vinyl alcohol), 275 277, 277f Rigid amorphous fraction (RAF), 229

400

S Semicrystalline biopolymer, DSC nonisothermal heating curve of, 109 110, 110f Semicrystalline polymers, glass transition of, 49 52, 50f, 51f Silk proteins, 107 108 DSC heating curves of, 109 110, 111f Specific heat capacity, of natural fibers, 117 119 experimental measurement, 119, 119f theoretical prediction, 117 119, 118f Spinline stress, 8, 10 12, 136 137, 139, 141 143 Spinneret, 8 10 Spinning, 8 13 dry, of poly(vinyl alcohol) fibers, 273 electrospinning of nylon-6, 237 239, 238f of nylon-6,6, 241 gel-spinning, 95 of poly(vinyl alcohol) fibers, 273 274 melt. See Melt spinning process, 2 4, 3f wet of poly(vinyl alcohol), 277 of poly(vinyl alcohol) fibers, 272 273 Stabilization, of poly(vinyl alcohol) fibers, 274 275 Suessen method, 235 Surface-to-volume ratio, TGA analysis of, 72 73, 72f Surgical sutures absorbable, 259 260 biodegradable, 260 definition of, 259 mechanical properties of, 259 nonabsorbable, 259 260 thermal analysis of, 259 T TA. See Thermal Analysis (TA) Temperature calibration of DSC, 23 36 on cooling, 26 30, 26f, 27f, 28f, 29f energy (enthalpy) calibration, 30 on heating, 23 25 specific heat capacity calibration (heat flow calibration), 31 36

Index

modulated temperature DSC, 36 traditional DSC, 32 35, 32f, 33f Temperature responsive fibrous materials, thermal analysis of, 335 future research, 350 phase-change material applications of, 347 350, 348f, 349f thermal management performance of, 342 347, 344f, 345f, 346f, 347f, 348f types of, 337 342 working principle of, 336 337, 337f TGA. See Thermogravimetric analysis (TGA) Thermal Analysis (TA), 4 5, 7, 81 of acrylic fibers, 297 of aliphatic nylon fibers, 223 of carbon fibers, 297 differential, 17 19 of liquid crystalline polymers, 325 of natural fibers, 105 of polyacrylonitrile fibers, 303 312 DMA characterization, 311 312, 313f DSC characterization, 303 312, 303f, 304f, 304t, 306f, 307t heating rates in air and nitrogen, TGA measurements of, 306 309, 307f, 309f, 309t PAN/IA/MA, PAN/IA, and PAN/VAc fibers, TGA measurements of, 309 311, 310t, 312f of polyester fibers, 139 145, 139f, 140f, 141f, 142f, 143f of polypropylene fibers, 211 219 of poly(vinyl alcohol) fibers, 271 Thermal insulation capacity of fireprotective fabrics, 361 368, 362f combined radiant-heat and gas flame insulation test, 364 365, 365f evaluation of, 374 382 insulation from combined radiant-heat and gas flame, 379 380 insulation from flash fire, 380 381 insulation from gas flame, 377 379 insulation from hot surface contact, 381 382 insulation from liquid flame, 381 insulation from radiant-heat, 375 377 flash fire insulation tests, 365 366, 366f

Index

gas flame insulation test, 363 364, 364f hot surface contact insulation test, 367 368, 368f liquid flame insulation test, 367, 367f radiant-heat insulation test, 361 363, 363f Thermal shrinkage, thermomechanical analysis of, 86 93, 88f, 89f, 90f, 91f, 92f Thermal stability fibers, differential scanning calorimetry of, 60 62, 61f of fire-protective fabrics, 358 361 evaluation of, 368 374 flammability test, 359 361, 360f transitional characteristics, 369 372, 370f, 371f, 372f transitional characteristics test, 358 359, 359f TGA analysis of chemical treatment effect, 78, 78f flame retardation effect, 78, 79f Thermal stress analysis (TSA) of nylon-6, 234 235 Thermogravimetric analysis (TGA), 71 of aliphatic nylon fibers, 230, 231f, 232f chemical treatment effect on thermal stability, 78, 78f crystallinity, 73, 74f flame retardation effect on thermal stability, 78, 79f of liquid crystalline polymers, 329 331, 330f, 331f of polyacrylonitrile fibers heating rates in air and nitrogen, 306 309, 307f, 309f, 309t PAN/IA/MA, PAN/IA, and PAN/VAc fibers, 309 311, 310t, 312f of polyetheretherketone, 251, 252f of polyetherketoneketone, 255 257, 256f, 257f polymer flammability, 73 75, 75f, 76f of poly(vinyl alcohol), 275, 276f radiation effect on pyrolysis, 75 77, 76f surface-to-volume ratio, 72 73, 72f

401

thermo-oxidative degradation, 73, 75f Thermomechanical analysis (TMA), 81 of aliphatic nylon fibers, 230 applications in fiber research, 82 83 coefficient of linear thermal expansion, 81, 84 86, 86f melt spinning of, 9 of nylon-6, 234 235 of poly(vinyl alcohol) fibers, 279, 282f shrinkage force, 86 93, 88f, 89f, 90f, 91f, 92f static force instrument, 83f Thermo-oxidative degradation, TGA analysis of, 73, 75f Tie molecules, 9 12 TMA. See Thermomechanical analysis (TMA) Transitional characteristics of fire-protective fabrics, 358 359, 359f evaluation of, 369 372, 370f, 371f, 372f U UHMWPE. See Ultra-high-molecular-mass polyethylene (UHMWPE) Ultra-high-molecular-mass polyethylene (UHMWPE), 198 199 V Vectran liquid crystalline film, DMA analysis of, 102, 103f W Wet spinning of poly(vinyl alcohol), 277 of poly(vinyl alcohol) fibers, 272 273 Wool fiber, chemical composition of, 370f Y Yarn after-processing, 14 Z Zone-annealing method, 235 236 Zone-drawing method, 235 236 “Zero-force” measuring technique, 81 82