Advanced green composites 9781119323266, 1119323266

Table of contents :
Content: Preface xiii 1 Introduction 1Anil N. Netravali 1.1 Introduction 2 2 Green Resins from Plant Sources and Strengthening Mechanisms 11Muhammad M. Rahman and Anil N. Netravali 2.1 Introduction 12 2.2 Green Resins from Agro-Resources 14 2.2.1 Plant Protein-Based Resins 14 2.2.2 Plant Starch-Based Resins 21 2.3 Green Resins from Microbial Fermentation 25 2.3.1 Polyhydroxyalkanoates 25 2.3.2 Pullulan 27 2.4 Green Resins Using Monomers from Agricultural Resources 29 2.4.1 Polylactic Acid 29 2.5 Strengthening of Green Resins using Nano-Fillers 32 2.5.1 Inorganic Nano-Fillers 33 2.5.2 Organic Nano-Fillers 38 2.6 Conclusions 43 References 44 3 High Strength Cellulosic Fibers from Liquid Crystalline Solutions 57Yuxiang Huang and Jonathan Y. Chen 3.1 Introduction 57 3.2 Fibers from Liquid Crystalline Solutions of Cellulose Derivatives 59 3.3 Fibers from Liquid Crystalline Solution of Nonderivatized Cellulose 60 3.4 Regenerated-Cellulose/CNT Composite Fibers with Ionic Liquids 61 3.5 Future Prospects 63 Summary 64 References 65 4 Cellulose Nanofibers: Electrospinning and Nanocellulose Self-Assemblies 67You-Lo Hsieh 4.1 Introduction 68 4.2 Electrospinning of Cellulose Solutions 70 4.3 Cellulose Nanofibers via Electrospinning and Hydrolysis of Cellulose Acetate 70 4.4 Bicomponent Hybrid and Porous Cellulose Nanofibers 72 4.5 Wholly Polysaccharide Cellulose/Chitin/Chitosan Hybrid Nanofibers 74 4.6 Surface-Active Cellulose Nanofibers 76 4.7 Nanocelluloses 77 4.8 Nanocelluloses from Agricultural By-Products 79 4.9 Source Effects - CNCs from Grape Skin, Tomato Peel, Rice Straw, Cotton Linter 80 4.10 Process Effect - Nanocelluloses from Single Source (Corn Cob, Rice Straw) 82 4.11 Ultra-Fine Cellulose Fibers from Electrospinning and Self-Assembled Nanocellulose 85 4.12 Further Notes on Nanocellulose Applications and Nanocomposites 87 Acknowledgement 88 References 88 5 Advanced Green Composites with High Strength and Toughness 97Anil N. Netravali 5.1 Introduction 98 5.2 'Greener' Composites 99 5.3 Fully 'Green' Composites 101 5.4 'Advanced Green Composites' 102 5.5 Conclusions 106 References 108 6 All-Cellulose (Cellulose-Cellulose) Green Composites 111Shuji Fujisawa, Tsuguyuki Saito and Akira Isogai 6.1 Introduction 111 6.1.1 Cellulose 111 6.1.2 Nanocelluloses for Polymer Composite Materials 112 6.1.3 All-Cellulose Composites 114 6.2 Preparation of ACCs 114 6.2.1 Dissolution of Cellulose 114 6.2.1.1 Aqueous Solvents 114 6.2.1.2 Organic Solvents 115 6.2.1.3 Ionic Liquids 115 6.2.2 Preparation of ACCs 116 6.2.2.1 One-Phase Preparation 116 6.2.2.2 Two-Phase Preparation 116 6.3 Structures and Properties of ACCs 120 6.3.1 Optical Properties 120 6.3.2 Mechanical Properties 120 6.3.3 Thermal Expansion Behavior 124 6.3.4 Gas Barrier Properties 124 6.3.5 Biodegradability 125 6.4 Future Prospects 125 6.5 Summary 126 6.6 Acknowledgements 127 References 127 7 Self-Healing Green Polymers and Composites 135Joo Ran Kim and Anil N. Netravali 7.1 Introduction 136 7.1.1 Self-Healing Property in Materials: What is it and Why it is Needed? 136 7.2 Types of Self-Healing Approaches Used in Thermoset Polymers 137 7.2.1 Microcapsule-Based Self-Healing System 138 7.2.1.1 Microencapsulation Techniques 139 7.2.1.2 Microcapsule Systems for Self-Healing 148 7.2.2 Vascular Self-Healing System 158 7.2.2.1 One-, Two-, or Three-Dimensional Microvascular Systems 159 7.2.3 Intrinsic Self-Healing System 161 7.2.3.1 Test Methods to Characterize Self-Healing 162 7.2.3.2 Quasi-Static Fracture Methods 163 7.2.3.3 Fatigue Fracture Methods 165 7.2.3.4 Impact Fracture Methods 166 7.2.3.5 Other Techniques 166 7.3 Self-Healing Polymers from Green Sources 167 7.3.1 Self-Healing Polymers in Biomaterials 168 7.3.2 Self-Healing Green Resins and Green Composites 170 7.4 Summary and Prospects 173 Acknowledgements 175 References 175 8 Transparent Green Composites 187Antonio Norio Nakagaito, Yukiko Ishikura and Hitoshi Takagi 8.1 Introduction 187 8.2 Cellulose Nanofiber-Based Composites and Papers 189 8.2.1 Bacterial Cellulose-Based Composites 189 8.2.2 CNF-Based Composites 191 8.2.3 Transparent Nanopapers 194 8.2.4 All Cellulose Transparent Composites 195 8.3 Chitin-Based Transparent Composites 197 8.3.1 Chitin Nanofiber-Based Composites 197 8.3.2 Micro-Sized Chitin Composites 199 8.3.3 Chitin-Chitosan Transparent Green Composites 200 8.3.4 All Chitin Nanofiber Transparent Films 202 8.4 Electronic Devices Based on CNF Films and Composites 202 8.5 Future Prospects 205 8.6 Summary 206 References 206 9 Toughened Green Composites: Improving Impact Properties 211Koichi Goda 9.1 Introduction 211 9.2 Significance of Fiber Length in Toughened Fibrous Composites 212 9.3 Impact Properties of Green Composites 217 9.3.1 Relation Between Interfacial and Mechanical Properties in Green Composites 217 9.3.2 A Pattern of Increase in Tensile Strength and Decrease in Impact Strength 221 9.3.3 Effect of Toughened Resin 227 9.3.4 Approaches to Increase Both TS and IS 228 9.4 Role of Large Elongation at Break in Regenerated Cellulose Fibers 229 9.5 Toughened Cellulose Fibers and Green Composites 231 9.5.1 Toughening Mechanism of Regenerated Cellulose Fibers 231 9.5.2 Mercerization Effect 234 9.5.3 Other Beneficial Chemical Treatments 238 9.6 Conclusions 240 Appendix 241 References 243 10 Cellulose Reinforced Green Foams 247Jasmina Obradovic, Carl Lange, Jan Gustafsson and Pedro Fardim 10.1 Introduction 248 10.2 Bio-Based Foams 249 10.2.1 Starch-Based Foams 250 10.2.2 Foams Based on Vegetable Oils 253 10.2.3 Foams Based on Poly(Lactic Acid) 255 10.3 Surface Engineering of Cellulose Fibres Used in Foams 256 10.3.1 Chemical Modifications of Cellulose Fibres 257 10.3.2 In Situ Synthesis of Hybrid Fibres 258 10.3.2.1 Topology and Particle Content on Hybrid Fibres 260 10.3.2.2 Foam Formation 262 10.3.2.3 Combustion Behavior of Foams 262 10.4 Prospects 265 10.5 Summary 266 Acknowledgements 267 References 267 11 Fire Retardants from Renewable Resources 275Zhiyu Xia, Weeradech Kiratitanavit, Shiran Yu, Jayant Kumar, Ravi Mosurkal and Ramaswamy Nagarajan 11.1 Introduction 276 11.2 Fire Retardant Additives Based on Phosphorus and Nitrogen from Renewable Resources 278 11.2.1 Nucleic Acids 279 11.2.2 Proteins Containing Phosphorus and Sulfur 286 11.2.3 Phosphorus/Nitrogen-Rich Carbohydrates 289 11.2.4 Carbohydrates 291 11.3 Natural Phenolic Compounds as Flame Retardant Additives 295 11.3.1 Lignin 296 11.3.2 Tannins 300 11.3.3 Cardanol and Polymers of Cardanol 306 11.3.4 Polydopamines 307 11.4 Other FR Materials from Renewable Sources 308 11.4.1 Chicken Eggshell 308 11.4.2 Banana Pseudostem Sap 308 11.5 Prospects 310 11.6 Summary 311 11.7 Acknowledgements 312 References 312 12 Green Composites with Excellent Barrier Properties 321Arvind Gupta, Akhilesh Kumar Pal, Rahul Patwa, Prodyut Dhar and Vimal Katiyar 12.1 Introduction 321 12.2 Biodegradable Polymers: Classifications and Challenges 323 12.2.1 Poly (lactic acid): Properties Evaluation, Modifications and its Applications 328 12.2.2 Cellulose Based Composites: Chemical Modifications, Property Evaluation, and Applications. 333 12.2.3 Chitosan Based Composites: Chemical Modifications, Properties Evaluation, and Applications 338 12.2.4 Natural Gum Based Composites: Chemical Modification, Property Evaluation and Applications 343 12.2.5 Silk Based Composites: Property Evaluation, Chemical Modifications and Applications 348 12.3 Summary 355 Acknowledgements 355 References 356 13 Nanocellulose-Based Composites in Biomedical Applications 369M. Osorio, A. Canas, R. Zuluaga, P. Ganan, I. Ortiz and C. Castro 13.1 Introduction 370 13.2 Nanocellulose Sources and Properties 370 13.2.1 Nanocellulose Sources 370 13.2.2 Nanocellulose Characteristics as Green Material 373 13.2.3 Nanocellulose Properties for Biomedical Composites 374 13.2.3.1 Mechanical Properties 374 13.2.3.2 Morphology 375 13.2.3.3 Surface Charge 375 13.2.3.4 Conformability 378 13.2.3.5 Thermal Properties 378 13.2.3.6 Non-Toxic 379 13.2.3.7 Biocompatibility 379 13.3 Biomedical Applications of Nanocellulose-Based Composites 379 13.3.1 Nanocellulose-Based Composites with Various Polymers 380 13.3.1.1 Polyvinyl Alcohol 380 13.3.1.2 Chitosan (Ch) 381 13.3.1.3 Acrylic Acid (AA) 382 13.3.1.4 Polyhydroxyalkanoates (PHAs) 382 13.3.1.5 Silk Fibroin 383 13.3.1.6 Polyaniline and Polypyrrole 383 13.3.1.7 Alginate 384 13.3.1.8 Collagen 384 13.3.2 Nanocellulose-Based Composites with Bioactive Ceramics 385 13.3.2.1 Hydroxyapatite (HA) 385 13.3.2.2 Iron Oxide Nanoparticles 385 13.3.2.3 Calcium Peroxide (CaO2) 386 13.3.2.4 Carbon Nanotubes 386 13.3.3 Nanocellulose-Based Composites with Metals 386 13.3.3.1 Silver Nanoparticles (Ag) 386 13.3.3.2 Gold Nanoparticles (Au) 387 13.4 Summary 387 13.5 Prospects 390 Acknowledgments 390 References 390 Index 403

Citation preview

Advanced Green Composites

Scrivener Publishing 100 Cummings Center, Suite 541J Beverly, MA 01915-6106 Publishers at Scrivener Martin Scrivener ([email protected]) Phillip Carmical ([email protected])

Advanced Green Composites

Edited by

Anil Netravali Department of Fiber Science & Apparel Design, Cornell University, Ithaca, NY, U.S.A.

This edition first published 2018 by John Wiley & Sons, Inc., 111 River Street, Hoboken, NJ 07030, USA and Scrivener Publishing LLC, 100 Cummings Center, Suite 541J, Beverly, MA 01915, USA © 2018 Scrivener Publishing LLC For more information about Scrivener publications please visit www.scrivenerpublishing.com. All rights reserved. No part of this publication may be reproduced, stored in a retrieval system, or transmitted, in any form or by any means, electronic, mechanical, photocopying, recording, or otherwise, except as permitted by law. Advice on how to obtain permission to reuse material from this title is available at http://www.wiley.com/go/permissions. Wiley Global Headquarters 111 River Street, Hoboken, NJ 07030, USA For details of our global editorial offices, customer services, and more information about Wiley products visit us at www.wiley.com. Limit of Liability/Disclaimer of Warranty While the publisher and authors have used their best efforts in preparing this work, they make no representations or warranties with respect to the accuracy or completeness of the contents of this work and specifically disclaim all warranties, including without limitation any implied warranties of merchantability or fitness for a particular purpose. No warranty may be created or extended by sales representatives, written sales materials, or promotional statements for this work. The fact that an organization, website, or product is referred to in this work as a citation and/or potential source of further information does not mean that the publisher and authors endorse the information or services the organization, website, or product may provide or recommendations it may make. This work is sold with the understanding that the publisher is not engaged in rendering professional services. The advice and strategies contained herein may not be suitable for your situation. You should consult with a specialist where appropriate. Neither the publisher nor authors shall be liable for any loss of profit or any other commercial damages, including but not limited to special, incidental, consequential, or other damages. Further, readers should be aware that websites listed in this work may have changed or disappeared between when this work was written and when it is read. Library of Congress Cataloging-in-Publication Data ISBN 978-1-119-32326-6 Cover image: Pixabay.Com Cover design by Russell Richardson Set in size of 11pt and Minion Pro by Exeter Premedia Services Private Ltd., Chennai, India Printed in the USA 10 9 8 7 6 5 4 3 2 1

Contents Preface 1

Introduction Anil N. Netravali 1.1 Introduction

2 Green Resins from Plant Sources and Strengthening Mechanisms Muhammad M. Rahman and Anil N. Netravali 2.1 Introduction 2.2 Green Resins from Agro-Resources 2.2.1 Plant Protein-Based Resins 2.2.2 Plant Starch-Based Resins 2.3 Green Resins from Microbial Fermentation 2.3.1 Polyhydroxyalkanoates 2.3.2 Pullulan 2.4 Green Resins Using Monomers from Agricultural Resources 2.4.1 Polylactic Acid 2.5 Strengthening of Green Resins using Nano-Fillers 2.5.1 Inorganic Nano-Fillers 2.5.2 Organic Nano-Fillers 2.6 Conclusions References 3

High Strength Cellulosic Fibers from Liquid Crystalline Solutions Yuxiang Huang and Jonathan Y. Chen 3.1 Introduction 3.2 Fibers from Liquid Crystalline Solutions of Cellulose Derivatives 3.3 Fibers from Liquid Crystalline Solution of Nonderivatized Cellulose

xiii 1 2 11 12 14 14 21 25 25 27 29 29 32 33 38 43 44 57 57 59 60 v

vi

Contents 3.4 Regenerated-Cellulose/CNT Composite Fibers with Ionic Liquids 3.5 Future Prospects Summary References

4 Cellulose Nanofibers: Electrospinning and Nanocellulose Self-Assemblies You-Lo Hsieh 4.1 Introduction 4.2 Electrospinning of Cellulose Solutions 4.3 Cellulose Nanofibers via Electrospinning and Hydrolysis of Cellulose Acetate 4.4 Bicomponent Hybrid and Porous Cellulose Nanofibers 4.5 Wholly Polysaccharide Cellulose/Chitin/Chitosan Hybrid Nanofibers 4.6 Surface-Active Cellulose Nanofibers 4.7 Nanocelluloses 4.8 Nanocelluloses from Agricultural By-Products 4.9 Source Effects – CNCs from Grape Skin, Tomato Peel, Rice Straw, Cotton Linter 4.10 Process Effect – Nanocelluloses from Single Source (Corn Cob, Rice Straw) 4.11 Ultra-Fine Cellulose Fibers from Electrospinning and Self-Assembled Nanocellulose 4.12 Further Notes on Nanocellulose Applications and Nanocomposites Acknowledgement References

61 63 64 65 67 68 70 70 72 74 76 77 79 80 82 85 87 88 88

5 Advanced Green Composites with High Strength and Toughness 97 Anil N. Netravali 5.1 Introduction 98 5.2 ‘Greener’ Composites 99 5.3 Fully ‘Green’ Composites 101 5.4 ‘Advanced Green Composites’ 102 5.5 Conclusions 106 References 108 6 All-Cellulose (Cellulose–Cellulose) Green Composites Shuji Fujisawa, Tsuguyuki Saito and Akira Isogai 6.1 Introduction

111 111

Contents vii

6.2

6.3

6.4 6.5 6.6

6.1.1 Cellulose 6.1.2 Nanocelluloses for Polymer Composite Materials 6.1.3 All-Cellulose Composites Preparation of ACCs 6.2.1 Dissolution of Cellulose 6.2.1.1 Aqueous Solvents 6.2.1.2 Organic Solvents 6.2.1.3 Ionic Liquids 6.2.2 Preparation of ACCs 6.2.2.1 One-Phase Preparation 6.2.2.2 Two-Phase Preparation Structures and Properties of ACCs 6.3.1 Optical Properties 6.3.2 Mechanical Properties 6.3.3 Thermal Expansion Behavior 6.3.4 Gas Barrier Properties 6.3.5 Biodegradability Future Prospects Summary Acknowledgements References

7 Self-Healing Green Polymers and Composites Joo Ran Kim and Anil N. Netravali 7.1 Introduction 7.1.1 Self-Healing Property in Materials: What is it and Why it is Needed? 7.2 Types of Self-Healing Approaches Used in Thermoset Polymers 7.2.1 Microcapsule-Based Self-Healing System 7.2.1.1 Microencapsulation Techniques 7.2.1.2 Microcapsule Systems for Self-Healing 7.2.2 Vascular Self-Healing System 7.2.2.1 One-, Two-, or Three-Dimensional Microvascular Systems 7.2.3 Intrinsic Self-Healing System 7.2.3.1 Test Methods to Characterize Self-Healing 7.2.3.2 Quasi-Static Fracture Methods 7.2.3.3 Fatigue Fracture Methods 7.2.3.4 Impact Fracture Methods

111 112 114 114 114 114 115 115 116 116 116 120 120 120 124 124 125 125 126 127 127 135 136 136 137 138 139 148 158 159 161 162 163 165 166

viii

Contents 7.2.3.5 Other Techniques 7.3 Self-Healing Polymers from Green Sources 7.3.1 Self-Healing Polymers in Biomaterials 7.3.2 Self-Healing Green Resins and Green Composites 7.4 Summary and Prospects Acknowledgements References

166 167 168 170 173 175 175

8 Transparent Green Composites Antonio Norio Nakagaito, Yukiko Ishikura and Hitoshi Takagi 8.1 Introduction 8.2 Cellulose Nanofiber-Based Composites and Papers 8.2.1 Bacterial Cellulose-Based Composites 8.2.2 CNF-Based Composites 8.2.3 Transparent Nanopapers 8.2.4 All Cellulose Transparent Composites 8.3 Chitin-Based Transparent Composites 8.3.1 Chitin Nanofiber-Based Composites 8.3.2 Micro-Sized Chitin Composites 8.3.3 Chitin-Chitosan Transparent Green Composites 8.3.4 All Chitin Nanofiber Transparent Films 8.4 Electronic Devices Based on CNF Films and Composites 8.5 Future Prospects 8.6 Summary References

187

9 Toughened Green Composites: Improving Impact Properties Koichi Goda 9.1 Introduction 9.2 Significance of Fiber Length in Toughened Fibrous Composites 9.3 Impact Properties of Green Composites 9.3.1 Relation Between Interfacial and Mechanical Properties in Green Composites 9.3.2 A Pattern of Increase in Tensile Strength and Decrease in Impact Strength 9.3.3 Effect of Toughened Resin 9.3.4 Approaches to Increase Both TS and IS 9.4 Role of Large Elongation at Break in Regenerated Cellulose Fibers

211

187 189 189 191 194 195 197 197 199 200 202 202 205 206 206

211 212 217 217 221 227 228 229

Contents ix 9.5 Toughened Cellulose Fibers and Green Composites 9.5.1 Toughening Mechanism of Regenerated Cellulose Fibers 9.5.2 Mercerization Effect 9.5.3 Other Beneficial Chemical Treatments 9.6 Conclusions Appendix References

231 231 234 238 240 241 243

10 Cellulose Reinforced Green Foams Jasmina Obradovic, Carl Lange, Jan Gustafsson and Pedro Fardim 10.1 Introduction 10.2 Bio-Based Foams 10.2.1 Starch-Based Foams 10.2.2 Foams Based on Vegetable Oils 10.2.3 Foams Based on Poly(Lactic Acid) 10.3 Surface Engineering of Cellulose Fibres Used in Foams 10.3.1 Chemical Modifications of Cellulose Fibres 10.3.2 In Situ Synthesis of Hybrid Fibres 10.3.2.1 Topology and Particle Content on Hybrid Fibres 10.3.2.2 Foam Formation 10.3.2.3 Combustion Behavior of Foams 10.4 Prospects 10.5 Summary Acknowledgements References

247

11 Fire Retardants from Renewable Resources Zhiyu Xia, Weeradech Kiratitanavit, Shiran Yu, Jayant Kumar, Ravi Mosurkal and Ramaswamy Nagarajan 11.1 Introduction 11.2 Fire Retardant Additives Based on Phosphorus and Nitrogen from Renewable Resources 11.2.1 Nucleic Acids 11.2.2 Proteins Containing Phosphorus and Sulfur 11.2.3 Phosphorus/Nitrogen-Rich Carbohydrates 11.2.4 Carbohydrates 11.3 Natural Phenolic Compounds as Flame Retardant Additives

275

248 249 250 253 255 256 257 258 260 262 262 265 266 267 267

276 278 279 286 289 291 295

x

Contents 11.3.1 Lignin 11.3.2 Tannins 11.3.3 Cardanol and Polymers of Cardanol 11.3.4 Polydopamines Other FR Materials from Renewable Sources 11.4.1 Chicken Eggshell 11.4.2 Banana Pseudostem Sap Prospects Summary Acknowledgements References

296 300 306 307 308 308 308 310 311 312 312

12 Green Composites with Excellent Barrier Properties Arvind Gupta, Akhilesh Kumar Pal, Rahul Patwa, Prodyut Dhar and Vimal Katiyar 12.1 Introduction 12.2 Biodegradable Polymers: Classifications and Challenges 12.2.1 Poly (lactic acid): Properties Evaluation, Modifications and its Applications 12.2.2 Cellulose Based Composites: Chemical Modifications, Property Evaluation, and Applications. 12.2.3 Chitosan Based Composites: Chemical Modifications, Properties Evaluation, and Applications 12.2.4 Natural Gum Based Composites: Chemical Modification, Property Evaluation and Applications 12.2.5 Silk Based Composites: Property Evaluation, Chemical Modifications and Applications 12.3 Summary Acknowledgements References

321

11.4

11.5 11.6 11.7

321 323 328

333

338

343 348 355 355 356

13 Nanocellulose-Based Composites in Biomedical Applications 369 M. Osorio, A. Cañas, R. Zuluaga, P. Gañán, I. Ortiz and C. Castro 13.1 Introduction 370 13.2 Nanocellulose Sources and Properties 370 13.2.1 Nanocellulose Sources 370 13.2.2 Nanocellulose Characteristics as Green Material 373

Contents xi 13.2.3 Nanocellulose Properties for Biomedical Composites 13.2.3.1 Mechanical Properties 13.2.3.2 Morphology 13.2.3.3 Surface Charge 13.2.3.4 Conformability 13.2.3.5 Thermal Properties 13.2.3.6 Non-Toxic 13.2.3.7 Biocompatibility 13.3 Biomedical Applications of Nanocellulose-Based Composites 13.3.1 Nanocellulose-Based Composites with Various Polymers 13.3.1.1 Polyvinyl Alcohol 13.3.1.2 Chitosan (Ch) 13.3.1.3 Acrylic Acid (AA) 13.3.1.4 Polyhydroxyalkanoates (PHAs) 13.3.1.5 Silk Fibroin 13.3.1.6 Polyaniline and Polypyrrole 13.3.1.7 Alginate 13.3.1.8 Collagen 13.3.2 Nanocellulose-Based Composites with Bioactive Ceramics 13.3.2.1 Hydroxyapatite (HA) 13.3.2.2 Iron Oxide Nanoparticles 13.3.2.3 Calcium Peroxide (CaO2) 13.3.2.4 Carbon Nanotubes 13.3.3 Nanocellulose-Based Composites with Metals 13.3.3.1 Silver Nanoparticles (Ag) 13.3.3.2 Gold Nanoparticles (Au) 13.4 Summary 13.5 Prospects Acknowledgments References Index

374 374 375 375 378 378 379 379 379 380 380 381 382 382 383 383 384 384 385 385 385 386 386 386 386 387 387 390 390 390 403

Preface Polymer composites are made of two distinct components: matrix or resin (continuous phase) and filler or reinforcement (discontinuous phase) and have properties that cannot be achieved by a single component alone. Conventionally, advanced composites have been defined as those that have excellent tensile properties. Three factors contribute to their high tensile properties: (1) high strength fibers such as carbon, aramid, or glass used as reinforcement, (2) good resins such as epoxies, and (3) excellent bonding between the fiber and the resin. Specific tensile properties of advanced composites are significantly higher than most metals because of their low density. As a result, advanced composites have replaced metals in many applications from aerospace to sports gears and from automobiles to wind turbines. It is being envisioned that future civil structures such as bridges and buildings will use advanced composites in place of steel, which will require significantly larger volumes. As we know, every material has a life span, and advanced composites are no exception. Unfortunately, we have not found an environmentally friendly way to dispose these composites at the end of their useful life nor have we found any sustainable raw material source to make them. Composites, at present, use unsustainable petroleum as the raw material to synthesize fibers and resins and end up in landfills at the end of their life, making that land useless for any other use for several decades or even centuries. Fortunately, this is slowly beginning to change with the advent of green polymers that are derived from fully sustainable plant based sources. Green composites are also being fabricated using these polymers with plant based fibers as reinforcement. At the end of their life, they can be easily composted, rather than being dumped into landfills. Thus, they can not only save the petroleum but also the land used for landfills. Significant research is going on around the world to develop green composites that use both fibers and resins that are derived from plants. However, composites using common fibers such as jute, kenaf, sisal, ramie, hemp, banana, and many others have only moderate tensile properties, comparable to wood and wood-based products such as particle xiii

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Preface

board and medium density fiber (MDF) boards. They are not suitable for structural applications where much higher strength and stiffness are required. Their durability, fire performance, and other functionalities are also not par with conventional composites. As a result, their applications have been very limited. In the last decade or so, however, many new developments in the area of green composites have come about that are changing the landscape drastically. For example, high strength cellulose fibers have been developed using liquid crystalline cellulose solutions and air-gap wet spinning technique used for spinning Kevlar fibers. Being 100% cellulose, they are fully degradable. Although this technology is still not mature, fibers with strength in the range of 1.5–1.7 GPa have been obtained. Furthermore, researchers have found ways to improve their properties further to around 2.0 GPa using chemical and heat treatments under tension. There is great promise that once the technology matures, these fibers could have tensile properties close to Kevlar . Advanced green composites with high strength and stiffness have already been fabricated using these fibers and soy protein and maize starch based resins that would be suitable for structural applications. Other than high strength and stiffness, composites are also being developed with a wide range of functional properties for a variety of applications. This has changed the conventional definition of Advanced Composites based on mechanical properties and now includes all composites that have special functional properties such as high transparency, fire resistance, ultra-light weight, autonomously repairing, and resin-less composites. Researchers have been able to obtain many of these properties in green composites as well, making them Advanced Green Composites. This book provides the current state of advanced green composites that have been developed or are at the research stage of being developed with a variety of functional properties. Chapter 1 presents a broad introduction to green composites and their development to date. In Chapter 2, Rahman and Netravali discuss green resins that have been derived from plant-based sources and seem to be promising to fabricate structural composites. The chapter also discusses some of the most promising bio-based and inorganic nano-fillers that are considered as potential candidates for enhancing the properties of these green resins. Improving resin tensile properties should automatically reflect on composite properties made using them. The third chapter by Huang and Chen discusses the development of high strength cellulosic fibers, the primary load-bearing component in current green composites. Viscose rayon process used for spinning cellulose fibers is more than a century old and is incapable of producing high strength cellulose fibers. This chapter provides an overview of high strength cellulosic

Preface

xv

fibers that can be obtained from liquid crystalline solutions of cellulose derivatives and nonderivatized cellulose as well as new methods to reinforce the liquid crystalline solutions of cellulose. It is to be noted that all high strength advanced green composites, until today, have been made using liquid crystalline cellulose (LCC) fibers. In the fourth chapter, Hsieh highlights top-down and bottom-up approaches to generate ultra-fine cellulose fibers of nano-scale dimensions, micro- and meso-porous and sheath-core hybrid structures as well as surface-functionalized fibrous materials. Micro- and nano-fibrillated cellulose (MFC/NFC) and other forms of nanocelluloses have high strength as well as aspect ratio and can be efficiently extracted from a variety of waste products such as apple, carrot, and orange pulps that remain after extracting juice, grape, and tomato skins, various straws, etc. They have been used as reinforcement in resins among many other applications. Chapter 5 discusses up to date efforts in developing high strength composites that use LCC fibers and are termed as advanced green composites. There are other opportunities for obtaining high strength composites as well. For example, researchers are trying to develop high strength fibers from spider-silk like protein. Once such fibers are commercially produced, they can be used to make green composites. Bacterial nanocellulose is another fiber with high strength. If these fibers can be oriented, they would be excellent as reinforcement as well. Chapter 6, by Fujisawa and colleagues, summarizes the latest in a new class of composites that do not require resin at all. These composites are made using only one component, cellulose, which acts as the reinforcing fibers as well as the resin that bonds the fibers. These all cellulose composites (ACC) are considered a green alternative to glass- and carbon-fiber-reinforced polymer composites. The authors also provide a future perspective on ACC development for applications in various fields, including optical devices, food, and medicine. Chapter 7 presents composites that have the ability to autonomously self-heal the damages such as microcracks, punctures, cuts, and scratches that result from the constant stress and strain they are subjected to during use. The damages continue to accumulate, ultimately failing the composites. Self-healing is designed to heal the damages as they occur and, hence, can increase the service life of the composites significantly. The chapter discusses different ways developed by researchers to achieve self-healing in conventional composites and how some methods have been extended to self-heal green resins and green composites. Selfhealing green resins and composites, with increased service life, should be more acceptable in mainstream applications in the future. In Chapter 8, Nakagaito and colleagues discuss optically transparent composites that have been developed of late for use in place of glass as substrates in

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Preface

electronic devices. Flexible electronics represents a common technology employed in gadgets that are ever-present in our daily lives. Among them, electronic displays are about to become flexible and foldable in the near future. Nanofibers that are invisible to our eyes can not only strengthen amorphous polymers but retain their transparency and also reduce their coefficient of thermal expansion, a critical requirement for electronic displays. Making transparent green composites is an active area of research at present and this chapter presents a thorough review of the current efforts. One of the deficiencies of the thermoset green resins such as plant-based proteins and starches as well as poly(lactic acid) is their brittle nature. That also translates into brittle composites. This is the topic of Chapter 9 in which Goda discusses ways of toughening composites with an emphasis on their impact properties. Chemical treatments of natural cellulosic fibers can make them stronger or tougher. Another key strategy to obtaining higher toughness is to control fiber/resin interfacial shear strength. Chapter 10 by Obradovic and colleagues introduces the science, technologies, and applications of ultralight porous green composites in the form of biofoams. The chapter presents a thorough review of the biofoam compositions, process methods, properties as well as their performance and applications. In Chapter 11, Xia and colleagues discuss fire retardants developed from renewable resources. Most fibers and composites used in construction, transportation, electronics, and protective textile applications must have fire retardant functionality. Commercially available halogenated flame retardants tend to be toxic and persist in the environment for a long time. The authors, based on the current state of research, conclude that exploring the use of renewable materials as feedstock for FR alternatives is quite promising. In addition to sustainability aspects, the research has unlocked exciting new possibilities in fundamental understanding of fire retardants and their mechanisms of action. Chapter 12 by Gupta et al. discusses the recent technological advancements and innovations in the development of novel biodegradable polymers and nanofillers derived from renewable feedstock with a strategy to convert “waste into wealth” with special focus on green composite films for stringent food packaging that require excellent barrier properties. Finally, in Chapter 13, Osorio and colleagues discuss nanocellulose-based composites in biomedical applications. Although cotton has been used in gauzes for treatment of wounds since ancient times, nanocellulose-based composites have become of great interest in biomedical applications, given their inherent biocompatibility. Nanocellulose from microorganism assembly is similar to that of collagen, the major component of extracellular matrices, and its applications are diverse, ranging from scaffolds for tissue engineering, implants for cell

Preface xvii regeneration or biosensors. Although the chapter discusses some relevant biomedical applications of nanocellulose based composites, authors predict that in medium to long term, nanocellulose-based composites will play an important role for developing in vitro tissues and organs, accelerating healing processes and improving life quality of mankind without impacting the environment. The book covering a broad range of green technologies should be of interest to researchers in academia, in government research labs, and R&D personnel in a host of industries (e.g., aerospace, automotive, biomedical, composites, fibers, medical, microelectronics, packaging, plastics, textiles, and others) who are interested in designing with green fibers, polymers, and composites or advancing the sustainable materials technology. Industries such as aerospace and automotive that are increasingly turning to composites for lightweighting each component but use conventional composites should be able to find greener alternatives in their applications. Anyone working in sustainable plastics/polymers and composites industries should find this book of great interest and very useful as well. It is my great pleasure to thank all those who made this book possible. First and foremost, I would like to profusely thank the outstanding authors who spent enormous amount of their valuable time in writing the chapters. Sharing their deep knowledge in the field and cutting edge research they are involved in, with the interested community, is greatly appreciated. This book would have been impossible to complete without their hard work, sustained interest, great enthusiasm, and cooperation. Thanks are also due to Martin Scrivener (Scrivener Publishing) for his unwavering support, interest, and encouragement, as well as patience in getting this book completed. Anil N. Netravali Cornell University Ithaca, NY [email protected] August 2018

1 Introduction Anil N. Netravali Department of Fiber Science & Apparel Design, Cornell University, Ithaca, NY, USA

Abstract Advanced composites using high strength fibers such as graphite and Kevlar have been replacing metals in load bearing structures for several decades. Initially introduced in aerospace applications where lighter structures are critical, the technology has become common and now being used in mundane applications such as automobiles and, in future, will be used in civil structures such as buildings and bridges. Obviously, there are several advantages of using these composites. However, disposal of these composites, at the end of their life, is already a problem and with large volume applications, e.g., automobiles and civil structures, it will be a significant issue. Also, the fibers and resins used in these composites are made from petroleum which itself is not sustainable. These issues along with government regulations all around the world on restricting the use of petroleum-based materials have created a great interest among the scientists to develop polymers, chemicals and composites using sustainable raw materials. Plants, grown every year, are perhaps the most sustainable resource for developing monomers and polymers. Also, cellulose, derived from plants and microbial sources, seems to be an excellent linear polymer for developing fibers. Composites using plant derived polymers (resins) and fibers, known as Green Composites, can already be found in the market. The next generation of green composites called advanced green composites are being developed that will have special functionalities such as high strength, stiffness or toughness, fully transparent, autonomously self-healing, fire retardant, nano-composites for biomedical applications, resin-less composites, light-weight foam composites, composites with gas barrier properties, etc. These will be able to compete with the petroleum based conventional composites in many applications. However, they will enjoy significant advantages such as composting at the end of their life rather than putting in landfills and never-ending raw material supply, i.e., fully sustainable. Rather than harming the environment, they will help it by

Corresponding author: [email protected] Anil Netravali (ed.) Advanced Green Composites, (1–10) © 2018 Scrivener Publishing LLC

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completing the nature’s intended carbon cycle. This book details the state-of-theart in advanced green composites.

1.1 Introduction Fiber reinforced composites based on fibers such as graphite, Kevlar , glass, etc., and resins such as epoxies, poly (etheretherketone) (PEEK), polyurethanes, etc., have been used in many structural applications from airplanes to windmill blades and from space stations to automobiles. These high performance composites possess extremely high strength and stiffness and when combined with their low density they exhibit high specific properties compared to all metals and alloys. As a result, they are commonly referred to as advanced composites. They also have other distinct advantages. They can be engineered to obtain desired properties by varying fiber forms, fiber placement and amount in different laminates. They can also be molded into desired shapes. In addition, they do not rust and can last for many decades when properly used. Their long-lasting nature, rather the ‘not-degrading’ characteristic, however, has created a real problem. Most plastics discarded on land or in sea, stay in ‘as is’ condition, i.e., without degrading, for many decades. This has not only created a severe litter problem but has also affected both wildlife and sea life in many ways. In case of composites, there is an additional issue that needs to be addressed. Because they contain two constituents, resin and fibers, that are bonded together, sometimes even covalently, at the end of their life they are impossible to be separated and reused. In case of thermoset resins such as epoxies, once cross-linked or cured, they cannot be re-formed or reused. At present, there is no acceptable environment friendly solution to dispose them off at the end of their life. And unfortunately, there is a life span for every composite, beyond which they cannot be used safely and, hence, have to be discarded. There are some efforts in reclaiming carbon fibers by burning off the resin. Some composites are incinerated to recover energy value. However, burning composites or resins creates toxic gases, which, if released into the environment, can lower the air quality. In addition, the burning process, along with additional handling involved, reduces the fiber strength significantly. Also, the fibers cannot be obtained in their original continuous form. For both these reasons, they can only be down cycled and may not be cost effective. Another way of recycling composites is to grind them into fine powder for use as filler in some applications. This also consumes a large amount of energy. At present, reclaiming fibers, recovering energy and grinding composites to powder are all carried out

Introduction 3 on a very small scale and can consume only a tiny fraction of the composites being discarded. As a result, most composites get discarded in landfills. Unfortunately, they stay in the landfills not just for several decades, but even centuries, without degrading. In fact, several studies have suggested that even the most biodegradable materials such as corn on the cob or even banana peel will not degrade in the anaerobic conditions of the landfills for decades. This can leave the land being used for landfills to be unavailable for any other productive use for centuries. Over the last three decades, the number of landfills in the US has gone down. As per the 1997 US census, the US had only 3091 active landfills. In addition, according to Leak Location Services, Inc., in 2000, 82% of the surveyed landfills had leaks which could pose threat to the water table. The tipping fees for landfills have also increased over the years. According to the Waste Business Journal, the average tipping fee in the US increased to $50.6 as of May, 2017, up by 3.5% from December 2016, in just 6 months! Since landfills are located far from the cities where most waste is generated, there is also a significant cost of transporting the waste to the landfills. While disposal is one critical problem associated with polymeric composites, their origin or the raw material used to derive them is the second big problem. Most commercial fibers, with the exception of glass fibers, and resins used in composites are produced using petroleum as the raw material. It is estimated that 5–7% of the total petroleum produced is used for making chemicals and materials such as polymers and composites. It is commonly accepted that petroleum will not last forever. We have been consuming petroleum at an unsustainable rate over the past 4–5 decades. By some estimates the consumption rate at present is 100,000 times faster than the earth can generate it. With countries such as China and India growing faster, petroleum consumption can only accelerate in the near future. At present to find new petroleum sources, we have to go deeper and farther into the sea. Many studies have suggested that at the current rate of consumption, we will have petroleum left for only 5–6 decades. That means in just over 100 years we will have consumed most petroleum earth took several million years create. As mentioned in the previous paragraph, at present, the only source to obtain plastics (polymers), fibers, and composites, is petroleum. So the questions we should be asking now are; what would happen when we run out of petroleum? From where will we get polymers/plastics that have become part of our daily life? How and which sustainable sources do we have to develop polymers, fibers and composites? While sustainable sources such as wind and sun (photovoltaic) exist for energy and are increasingly being used, the only sustainable source for obtaining materials is Plants.

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The growing global environmental consciousness about the persisting litter problem, concerns regarding air, water and land pollution, high rate of depletion of petroleum resources have resulted in new environmental regulations all over the world. Many governments have put in place laws to curtail the use of petroleum based materials and moved to sustainable plant-based materials that are also biodegradable and, hence, can be easily composted. The laws are also in place to collect and recycle polymers/ plastics. All these factors have together triggered the search for new products and processes that are environment friendly. Companies have realized that their existence will be in jeopardy in just 3–4 decades if they do not find a suitable substitute for petroleum that is sustainable. ‘Green’ chemistry, ‘green’ materials, recycling, sustainability, cradle-to-cradle design, life cycle analysis, industrial ecology, are no more ‘feel good’ buzz words but have begun guiding the developments of many new products. Researchers from academia, industry and governments world over have seriously taken the challenge in exploring plants and agricultural/food processing wastes as sources for chemicals and polymers/plastics and to design new products. This realization alone is significant and progress is being made in this direction. For example, plant-based polylactic acid (PLA) has already been a commercial success. Ethylene glycol developed from plant source is now being used for making poly (ethylene terephthalate) (PET) used in soda bottles and apparel. Polyethylene (PE) and polypropylene (PP) can also be made from sugar obtained from sugarcane. Poly (butylene succinate) (PBS) is increasingly made from plant source. Research has even indicated that resins such as epoxies can also be made from sugars. Many more examples of other polymers can be found at research stage and will be commercialized in the coming years. The above discussion should convince anyone that the materials world is experiencing a true Green Revolution, perhaps because of the realization that there is no other alternative. Fiber reinforced composites are no exception to this new paradigm. Past 2–3 decades have seen exponential rise in research in the field of Green Composites. This is clear if the number of papers and books published and research presented at conferences is any indication. Since plant based cellulosic fibers have been available and used universally for centuries, it has been easy to combine them with various resins as reinforcing constituents to fabricate composites. They are also available in a variety of forms, from fibers to yarns and from fabrics to nonwoven mats. This makes it easy to combine various forms to engineer the properties of composite as desired for the applications. Many examples can be found where plant fibers have been combined with resins such as epoxies, PE, PP, PET, PVC, etc., forming what are termed as ‘Greener

Introduction 5 Composites’. These greener composites have properties comparable to or even better than wood and wood based products such as plywood, particle boards, medium density fiber (MDF) boards, etc. They have been successfully used in applications such as automobile dashboards and door panels, ceiling tiles, wall panels, furniture, shelving, z-truss structures for flooring in recreation vehicles, etc. Since these greener composites combine degradable fibers and non-degradable resins, they cannot return to industrial metabolism or natural metabolism. They can only be down cycled to make low value products, incinerated to recover energy value or put in the landfills. The current trend is to fabricate composites that are fully ‘Green’ by using both resins and fibers that are plant based, sustainable, biodegradable and using processes that are water based, where possible, in effect, reducing their carbon footprint to zero. In some cases, it may be even possible to obtain a negative carbon footprint by preserving the carbon sequestered by nature. Many resins such as those based on plant proteins, starches, other gums, as well as synthesized biodegradable polymers such as PLA, have been used to fabricate green composites. While still in its infancy, research in this area has been on the rise in the past 2 decades. At the end of their life green composites can be composted without harming the environment. In fact, composting can create organic soil that can be used in growing more plants, thus, helping the environment…a truly winwin situation! Green composites can be used in many applications such as mass-produced consumer items that are used only once or just a few times before discarding. As many of these green composites are made using resins based on plant proteins or starches, they tend to be hydrophilic and absorb moisture when exposed to humid conditions or water (rain) and swell. This swelling is reduced when the composites are placed in dry environment and desorb the moisture. Cyclic absorption/desorption of moisture invariably degrades the fiber/resin bond that is accompanied by loss in mechanical properties. While one way to solve this issue is to limit these hydrophilic composites to indoor applications, another way would be to use paints, varnishes or hydrophobic coatings, as is commonly done to wooden articles. The latter method can allow using green composites in outdoor applications. Many examples can be found of green composites made using natural cellulosic fibers. However, since the natural cellulosic fibers have average strength of around 400 MPa, these composites cannot be strong. Most of these composites have strength in the range of 150 MPa to 250 MPa and can be used as replacement for wood, particle boards, MDF boards in applications such as cabinetry, door panels, etc., that do not require high

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strength. These composites, as mentioned earlier, are suitable for indoor applications. For outdoor applications, they need to be protected by hydrophobic paints, varnishes, or coatings. By conventional definition, composites with high strength and/or high modulus such as those made using glass, graphite or Kevlar fibers are called as Advanced Composites. They are commonly used in load bearing structural applications in place of metals to reduce weight. In the same spirit green composites that possess high strength, modulus and/or toughness have been called Advanced Green Composites. Such composites, up until now, have been successfully made using liquid crystalline cellulose (LCC) fibers that have been recently developed at Groningen University in The Netherlands. LCC fibers have higher crystallinity and molecular orientation than any other continuous cellulose fibers, natural or regenerated. In addition, their original strength of about 1.5 GPa has been successfully increased to about 2.0 GPa by combined chemical and mechanical treatments carried out at temperatures of 50 to 80 oC. The same combination of treatments has also increased their modulus from about 45 GPa (original) to close to 70 GPa making them a prime reinforcing constituent to fabricate advanced green composites by combining them with green resins. Such composites, made with hand layup, resulted in strengths of around 800 MPa and modulus of over 30 GPa with just 55% LCC fiber volume. With fiber volume between 65 and 70%, normal for any fiber reinforced composites, and eliminating some of the defects created by hand layup, the strength of these advanced green composites could be increased to over 1 GPa and modulus to over 40 GPa. Since these fibers have higher fracture strain compared to conventional fibers such as graphite, glass or Kevlar , they can absorb significant amount of energy during fracture. Such advanced green composites could be suitable not only for load bearing structures but also in ballistic applications. Another significant advantage of the plant based cellulosic fibers lies in their fibrillar nature. While the fibrils themselves may not be oriented along the fiber axis, the molecular orientation and crystallinity within the fibrils is very high. As a result, cellulose nanofibrils (CNFs), microfibrillated cellulose (MFC) or nanocrystals (CNCs) obtained from plant fibers have excellent tensile properties and have been used to strengthen and stiffen protein and starch based resins. CNFs have also been shown to improve resin toughness by crack-bridging mechanism. The modulus and strength of CNFs has been estimated at 140 GPa, and between 2 and 6 GPa, respectively. It should be noted that these values are close to or better than those of Kevlar . In addition to being biodegradable and sustainable, CNFs and CNCs can be obtained from waste products such as used newspaper,

Introduction 7 and apple, orange or carrot pulps that remain after extracting juice and are often discarded as waste. It has also been found that adding CNFs or MFC, with high aspect ratio, to green resins can improve the fiber/resin mechanical bonding in composites by providing possibilities for entanglement between the two, thus, further improving the composite properties. There are also abundant opportunities to use bacterial nanocellulose (BC) to obtain advanced green composites with high mechanical properties. BC nanofibers are 50–70 nm in diameter or width. BC is produced by aerobic bacteria such as acetobacter xylinum and has cellulose-I crystal structure with extended chain conformation. However, compared to plant based cellulose, BC has a higher degree of polymerization. It is also extremely pure form of cellulose and does not contain hemicellulose, lignin, pectin or wax that plant based cellulose fibers contain. As a result, BC nanofibers have high strength and stiffness and can be useful for fabricating advanced green composites. However, since bacteria move in a random fashion in the culture medium, most BC pellicles obtained from conventional process contain nanofibers that are randomly oriented and highly entangled. It is well known that in a composite, randomly oriented fibers cannot contribute to the strength and stiffness fully. Most advanced composites contain fibers that are highly oriented within individual laminates. As a result, there have been some efforts in orienting the BC and to take advantage of their high strength. Efforts using poly (dimethylsiloxane) (PDMS) substrates with micro-channels and restricting the bacterial movement inside those channels to form BC arrays and drawing them further to improve the orientation have been successful. This process has resulted in significant improvement in BC orientation which, in turn, resulted in enhanced strength and stiffness of the starch based green composites reinforced with them. Never the less, the results have also indicated that significant scope exists to improve the BC orientation further by optimizing the channel design and substrate geometry. Other methods to obtain highly oriented BC may be developed that are easier than using PDMS substrates with channels and may result in higher productivity. Besides obtaining higher mechanical properties, in the past couple of decades, researchers have achieved significant progress in enhancing various functional properties of green composites, making them ‘advanced’ in many ways. Functionalization of composites has clearly broadened the definition of Advanced Green Composites to encompass all green composites that have special functional capabilities. This book presents advanced green composites that have been developed with many functional properties. For example, all-cellulose composites that do not have any resin. The cellulose fibers used in these composites are bonded not by

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any resin but by cellulose itself. This is achieved by dissolving a very thin outer layer of cellulose fibers by a suitable solvent. The dissolved cellulose, in amorphous state, allows bonding of neighboring fibers. Once the fibers are bonded, the solvent is removed. All-cellulose composites can be made in the form of membranes that are strong and, hence, suitable for many applications. Another functional property is transparency. Fully transparent green composites have been made using cellulose CNFs, CNTs, or BC that have diameters less than 100 nm, much lower than the wavelength of visible light, and transparent green resin such as amorphous PLA. The CNFs, CNTs, and BC not only strengthen the composites but toughen them as well because of their crack bridging action. Transparent green composites can have many applications in light-weighting products by replacing glass which tends to be not just heavier but brittle as well. Green composites with brittle resins such as soy protein and starch are known to develop microcracks when subjected to stress. These microcracks slowly grow with time and all that is needed is for one of them to reach catastrophic level and fracture the composite. The catastrophic growth of the microcracks can be prevented if the microcracks can be healed as and when they are formed. This is commonly accomplished by releasing a healant which stays encapsulated in either microcapsules or a network of vascular system. Self-healing green resins, protein and starch based, and composites have now been developed that can autonomously heal in a similar fashion compared to conventional composites. However, all materials used in the case of green composites are green and sustainable. The healing mechanism in this case is obtained by fracture of the microcapsules that come in the path of the microcracks, releasing the healant and healing them. This is a new development and can keep the green composites operating for a much longer time, thus, extending their life significantly and making them economical than the current ones. Light weight reinforced green composite foams have also been made with the help of reinforcing CNFs and CNTs. These reinforced foams can be used for insulation or packaging. Much research has been conducted on synthesizing green fire retardant chemicals and fabricating fire resistant green composites. Proteins, used as resin in green composites, do not burn easily because of their inherent chemical composition that contains large amount of nitrogen. In addition, they also contain polar groups such as hydroxyl, carboxyl and amine groups. In normal environment, these polar groups absorb water which can help retard or extinguish the fire. In addition, when heated, there are possibilities that hydroxyl and amine groups could react with adjacent carboxyl groups to form ester and amide

Introduction 9 groups, respectively. These reactions not only cross-link the protein and tighten the structure but every time these condensation reactions occur, they eliminate a water molecule which can help retard the fire. In addition, the intumescent char produced by the fire on the composite surface can also significantly reduce the fire effects. Protein based green composites tend to be inherently fire retardant. Green composite films and membranes with excellent barrier properties have also been developed for stringent food packaging. These have high gas barrier properties and can preserve foods equally well compared to conventional petroleum based polymers. They are made using an existing wide pool of bionanofillers derived from biomass. And like all green composites, these films can be easily composted. Nanocellulose, as stated earlier, is also biocompatible. Nanocellulose can be used to mimic many tissues such as cartilage, skin, blood vessels, etc. Nanocellulose based biomedical green composites are predicted to play an important role in developing in vitro tissues and organs, accelerating healing and improving overall quality of life. In summary, great progress has been made in the field of green composites in just the past 2–3 decades. Many special functional characteristics, apart from high strength and stiffness, have been achieved in these composites making all of them advanced green composites. The field of advanced green composites is still in its infancy but with significant research going on all over the world, it won’t be much longer before these composites are commercialized and give conventional advanced composites a run for their money.

2 Green Resins from Plant Sources and Strengthening Mechanisms Muhammad M. Rahman and Anil N. Netravali* Department of Fiber Science & Apparel Design, Cornell University, Ithaca, NY, USA

Abstract Polymers derived from fossil resources have been the most widely used materials for applications ranging from packaging to sporting goods and from structural to aerospace applications in the past few decades. This is primarily due to their low price, light weight, uniform properties and matured production technology. However, increased concerns associated with the environmental deterioration, waste disposal difficulties and costs, and fossil resource depletion have fueled a growing interest in the development of renewable, environment-friendly and biodegradable polymers with comparable performances. Considerable interest has already been generated in the development of ‘green’ composites based on sustainable and biodegradable polymeric resins such as proteins, starch, etc. and fibers such as cellulose. Consequently, a few of the plant-based green polymeric resins have already replaced some non-degradable synthetic resins. Low-density combined with excellent mechanical properties and full biodegradability of these green composites compared to synthetic composites can provide lighter and stronger eco-friendly structures. This chapter focuses on the important constituent of green composites, biodegradable green resins that have proven to be promising to fabricate structural composites. The chapter also discusses some of the most promising bio-based and inorganic nano-fillers which are considered as potential candidates for enhancing the properties of these green resins which can result in better composite properties. Keywords: Bio-based resins, strengthening of resins, green resins, green composites, mechanical properties

*Corresponding author: [email protected] Anil Netravali (ed.) Advanced Green Composites, (11–56) © 2018 Scrivener Publishing LLC

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2.1 Introduction Composites are engineered materials that are fabricated using two or more constituents with significantly dissimilar mechanical and/or other functional properties to achieve properties unattainable from either of the constituents alone. Typically, in a polymeric composite, the constituents are called resin (continuous phase) and particulate fillers or fibrous reinforcement (dispersed phase). The resin acts as a binder for the fiber or particulate fillers in polymeric composites. While the fibers act as the principal load carrying member in composites, the resin acts as the load transfer medium between them and keeps the fibers in desired location and orientation. Particles in composites may sometimes act as load carrying member or resin stiffening member, depending on their size, shape and dispersion. More often than not, however, they act as simple non-reinforcing fillers. Most constituents that make conventional composites are derived from petroleum. Such composites being designed for long term applications, do not degrade for several decades or even centuries in normal environment and, as a result, have been responsible, at least partially, for the significant land and water pollution that exists at present. Because the composites cannot be easily recycled or reused, most of them are discarded in landfills at the end of their useful life. Increased concerns associated with the environmental deterioration, waste disposal difficulties and costs, and rapid depletion of petroleum source have fueled a growing interest in the development of ‘green’ resins from sustainable sources for fabricating high performance composites that can replace petroleum based conventional composites [1–3]. Green resins are considered to be those resins that can decompose completely in natural aerobic and anaerobic conditions without harming the environment [4]. In the past two decades, development of green resins and composites has received considerable attention due to environmental issues and economic concerns [5, 6]. As a consequence, a few of the green resins have already replaced some non-degradable synthetic resins in various structural applications [2, 4]. Most of these resins are derived from renewable resources such as plant- and animal-based raw materials. This chapter focuses on the largest category of resins that are derived from plant-based and fully sustainable raw materials as well as those derived from microbial sources. These resins can be categorized into three groups based on their origin and production. The first group includes resins that can be obtained directly from various agro-resources. Typical examples of these resins include polysaccharides, proteins, lignin, lipids, natural rubber, etc. The second group of green resins consists of those that are derived from microbial fermentation. For example, polyhydroxyalkanoates

Green Resins and Strengthening Mechanisms 13 (PHAs) and their copolymers have been synthesized by a wide variety of Gram-positive and Gram-negative bacteria cultivated using agricultural raw materials. The third group consists of resins that have been chemically synthesized using monomers from renewable resources such as lactic acid and/or an assortment of amino acids. Some prime examples that fall into this category are polylactic acid (PLA), polyaspartic acid, and polyglutamic acid, etc. These categories along with some typical examples are presented in Figure 2.1. The biodegradation of these resins can be achieved primarily through the enzymatic action of microorganisms such as bacteria, fungi, and algae present in the environment. These microorganisms can metabolize the resin to produce CO2, CH4, water, biomass, humic matter and other natural materials that get easily blended with natural soils and act as organic fertilizers and/or soil conditioners [4]. These resins, however, can be degraded by nonenzymatic processes such as chemical hydrolysis as well [4]. Hence, green resins are naturally recyclable and renewable through biological processes. Among various biodegradable polymers investigated so far for various technical applications, this chapter focuses on plant-based polymers such as proteins, starches, pullulan, and PLA as well as those derived from microbial sources such as PHA, that are most widely accepted and have proven to be promising green resins to fabricate fiber reinforced composites. While some of these green resins have been highly attractive for use as an alternative to some petroleum-based resins in terms of availability and cost, their mechanical properties limit their use in load-bearing structural

Green resins from plant sources

Resins directly extractd from plants Protein Edible protein (soy, wheat, pea, etc.) No-edible protein (jatropha, castor, etc.)

Resins derived from microbial fermentation Polyhydroxyalkanoates (PHAs) Pullulan

Starch Edible starch (corn, potato, etc.) No-edible starch (mango etc.)

Figure 2.1 Various categories of green resins.

Resins synthesized chemically from plant-based monomers Polyactic acid (PLA)

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applications. Various nano-fillers have shown the ability to improve the mechanical properties of these resins, due to their high specific surface area, high strength and stiffness. This chapter discusses some of the most promising bio-based nano-fillers which are considered as potential candidates for enhancing the properties of green resins. Green resins with higher properties can then be used to fabricate green composites.

2.2

Green Resins from Agro-Resources

2.2.1 Plant Protein-Based Resins Among the myriad of renewable resources available from plants, natural proteins obtained from seeds and beans have attracted significant interest as a new class of green resin for fabricating eco-friendly and biodegradable green composites because of their desirable functional properties that allow chemical modifications [7–9]. Proteins, also known as polypeptides, are one of the most essential macromolecules in biological systems. The basic monomeric units of all proteins are amino acids. The amino acids are connected by amide (peptide) bonds (-NHCO-) formed by the reaction between amine (-NH2) and carboxyl (-COOH) groups that are present in the amino acids. Figure  2.2 shows the protein primary structure, amino acid residues and peptide bonds. Proteins are highly complex polymers composed of up to twenty different amino acids. The amino acids have been divided into several categories such as polar, non-polar, aromatic, anionic and cationic based on their side groups [9]. The protein structure is organized into four levels: primary, secondary, tertiary, and quaternary structures. Figure  2.3 shows the four levels of protein structure.

Polypeptide chain

H

Peptide bond

Amino acids

N

Amino acids Phe Leu Ser Cys

O

H

H

C

H

C O

R

Figure 2.2 Protein primary structure, amino acid residues and peptide bonds [11].

Green Resins and Strengthening Mechanisms 15 The primary structure of proteins is a linear sequence of amino acids that are bonded covalently by peptide bonds while the secondary structure has two major forms, namely, α-helices (helical conformation) and β-pleated sheets (sheet-like conformation) because of the various interactions among the amino acids such as hydrogen, ionic or van der Waals interactions [9]. The folding of secondary structure by nonlocal interactions, salt bridges, hydrogen bonds, and disulfide bonds between thiol (sulphydryl) groups (–SH) from cysteine units forms a 3-D tertiary structure which can be assembled into more complex form by non-covalent bonding is termed quaternary structure [9]. The complex structure of proteins through noncovalent bonds can be strongly affected by factors such as temperature, pH, and organic solvents. The changes in quaternary, tertiary and secondary protein structures by the application of such factors as heat, mechanical agitation, and high or low pH are designated as protein denaturation. Another property related to protein structure, related to the presence of basic and acidic groups in their amino acids, is termed as isoelectric point. The isoelectric point of proteins is defined as the pH at which the protein has no net charge, i.e., the number of positively charged residues is equal to the number of negatively charged ones [9]. The protein molecules aggregate at their isoelectric point and become insoluble in water [10]. Below or above the pH of isoelectric point, the protein has a net positive or negative charge which facilitates the denaturation (opening) of the proteins. Under such conditions, i.e., away from the isoelectric point, the molecules open up and can dissolve in water. Opening up of the molecules also exposes the polar groups and makes it possible to react with other chemicals such as cross-linkers. The extraction process of proteins from protein-rich plant seeds is very straight-forward, under laboratory conditions. At first, the oil or fat (commonly, 20 to 40% of the seed weight) is extracted by either solvent extraction or screw-pressing method. Usually, solvent extraction is more efficient method compared to mechanical screw-pressing in terms of fat extraction from seeds. The residual defatted seed cakes are then ground into powder and sieved to obtain homogeneous micro particles. The sieved powder is

Pleated sheet

Pleated sheet

Alpha helix

Alpha helix

Amino acids

Primary structure

Secondary structure

Tertiary structure

Figure 2.3 Levels of protein structure organization [12].

Quaternary structure

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then suspended in deionized (DI) water and adjusted to alkaline pH. The insoluble residue from the suspension is discarded after stirring and the supernatant is collected. The pH of the supernatant is then adjusted to acidic pH and stirred for isoelectric precipitation of protein. The resulting precipitation is separated by centrifugation and washed using an ethanolwater mixture. The washed protein is then lyophilized, milled and stored for multiple applications including preparation of protein-based resin. The protein-based resin can be processed by two different methods, namely ‘wet method’ and ‘thermoplastic method’ [13]. One of the easiest and least expensive ways to prepare the protein-based resins is the wet method. At first, protein powder is homogenized in a suitable solvent such as water and ethanol and various plasticizers such as polyols (glycerol, polyethylene (PE) glycol, triethylene glycol, diethylene glycol, ethylene glycol), sorbitol, and sucrose are added [14]. The addition of plasticizers is important to make tougher resin by reducing the brittleness as the small plasticizer molecules can add significant amount of free volume to the system and increase the protein chain mobility. The protein in the solution is denatured, commonly, by adjusting both pH and high temperature and also by mechanical stirring. At this stage, the resin is termed as ‘pre-cured’ resin. The pre-cured resin solution may be dried and hot-pressed to complete the curing for further characterization. To improve the mechanical, thermal and water-resistance properties, various cross-linkers such as glyoxal, glutaraldehyde, formaldehyde, ferulic acid, transglutaminase, and sugar aldehydes, etc., have been used to cross-link the plasticized protein-based resin solution during the pre-curing stage [15, 16]. Such cross-linked proteins form thermoset resins that are suitable for fabricating green composites. To prepare thermoplastic resin, proteins are processed above their glass transition temperature which is commonly reduced with the help of plasticizers. Most processing is carried out at higher temperatures which convert the material into a rubbery mass which can be shaped into the desired shape and then stabilized by cooling or by evaporating the volatile plasticizers. Thermoplastic processes such as compression molding and extrusion have been popular in industry since they can produce films and coatings on a larger scale as well as in a very short time, resulting in high productivity. Some studies have been conducted for improving the sheet properties via cross-linking induced by heat [17, 18], enzymes [18], or irradiation [17]. Plant proteins, mainly obtained from edible oilseeds and beans/grains of various crops, have been extensively used as resins in green composites because of their abundant production and availability in the desired form throughout the world as well as their inexpensiveness [5, 7, 8, 16]. Soybean protein is one of the most important as well as the most widely

Green Resins and Strengthening Mechanisms 17 available edible protein [19]. It has been considered as one of the major green resins that can be processed into green composites using modifiers such as plasticizers, cross-linkers, fillers, and reinforcing agents [20–26]. Soybeans contain 40–45% protein, 18–20% fat, and 25–30% carbohydrates on a dry weight basis. Commercially, soy protein is available in three different forms, namely, soy flour (SF, 53–55% protein), soy protein concentrate (SPC, 65−72% protein), and soy protein isolate (SPI, ≥90% protein) on a dry weight basis [24]. Pure soy protein has molecular weights ranging from 140 to 360 kDa, consisting of different types of fractions, namely 2S, 7S, 11S, 15S fractions where S stands for Svedberg units [9]. Two major fractions in soy protein are 7S (35% of total soy protein and a molecular weight of about 150–190 kDa) and 11S (52% of total soy protein and a molecular weight of 320–360 kDa) [9]. These fractions have various polar functional groups in their amino acids, such as carboxyl, amine and hydroxyl groups which are capable of chemically reacting to form cross-links and improve the mechanical and physical properties of the proteins as well as their water resistance. Like soy protein, other edible proteins such as wheat gluten, corn zein, peanut and pea proteins have also been considered as eco-friendly resins for the production of composite materials for a broad spectrum of applications since they have excellent film-forming and adhesive capabilities [13, 15, 16, 27, 28]. Peanut seed contains about 45% fat and 22 to 33% protein while defatted peanut has 55 to 60% protein [29]. It is also an important agricultural product in the United States. Wheat gluten is the storage protein of wheat. Wheat kernel has 8 to 15 wt% wheat protein. The water insoluble fractions of wheat proteins are prolamin and glutelin. The prolamin fraction is also called gliadin while the glutelin fraction is also termed as glutenin. Both represent about 33% and 46% of the total protein content in wheat, respectively [15]. Zein is the major storage protein in the corn endosperm which accounts for 45 to 50% of the dry weight [30]. It is an alcohol-soluble protein with a molecular weight of 18 to 45 kDa. The poor solubility of zein in water is due to its high content of hydrophobic amino acid residues. Pea is a pulse crop whose annual production is 10–12 million tons in the world [31]. The protein content of dry peas ranges from 14 to 39 wt% and contains mainly globulin, which accounts for 65–80% of the protein [31]. The globulin fraction contains legumin (MW 360–400 kDa), vicilin (MW 160–200 kDa), and small quantities of convicilin (MW 280 kDa) [31]. In general, protein films show excellent barrier properties to oxygen, lipids, and aromas but have only moderate mechanical properties. However, they have high water permeability due to their hydrophilic character. The mechanical, thermal and physical properties of various edible protein-based resins are given in Table 2.1.

140–300

30–300

21–25

25–50

160–400

Soy [32–34]

Wheat gluten [35,36]

Zein [31]

Peanut [29,37,38]

Pea [31,39,40]

85.2–251.3

0.9–704

49.9–1967

1.2–811.9

61–1667

Young’s modulus (MPa)

0.5–7.5

0.6–12.3

1.6–40.6

0.8–22.2

1.0–40.6

Tensile strength (MPa)

30.6–150.0

3.9–160

0.8–165

8.5–301

3–175

Fracture strain (%)

NA

NA

164–168

162

145

Tg** ( °C)

82.1

NA

270

292–293

250–275

Td*** ( °C)

* The properties show a broad range because of the modifiers used. ** Tg is the glass transition temperature at which the transition between the glassy and rubbery state occurs in the amorphous regions of a polymer. *** Td is the thermal degradation temperature at which the polymer chemically decomposes.

Molecular weight (kDa)

Protein

Table 2.1 Properties of various edible protein sheets*.

18 Advanced Green Composites

Green Resins and Strengthening Mechanisms 19 Although edible proteins have been, thus far, used extensively to fabricate ‘green’ resins, there have been ethical questions raised on the utilization of these food sources for technical applications on a large scale. As a result, there have been a few studies that have focused on exploring nonedible protein sources for various technical applications so as to resolve the issue of food security. Proteins from some oil-rich seeds such as neem (Azadirachta indica), karanja (Pongamia pinnata), castor bean (Ricinus communis L.), sunflower (Helianthus annuus), and jatropha (Jatropha curcas) which are considered to be non-edible because of the presence of toxic and anti-nutritional compounds to human and animals are now being researched for various technical applications [41–48]. They do not compete with the edible protein sources and can be used as resins in composites without raising any ethical questions. In recent years, many of these seeds, particularly jatropha, have been utilized as a source of biodiesel production [49, 50]. After extraction of oil from the seeds for biodiesel, protein-rich defatted seed cakes remain as waste residues and may be utilized as very low-value applications such as organic nitrogenous fertilizers or fuel sources [43, 51]. Another inherent advantage of using these proteins is that the presence of toxic compounds can altogether eliminate the need for antimicrobial and anti-fungal compounds, that are commonly needed for the edible proteins to make them safe from the attacks of microbes or insects for an extended period [43, 46]. Neem seeds, used for medicinal purpose and pest management contain 40 to 50% oil and 40 to 50% protein [45]. Another plant ‘karanja’, used for biofuel extraction, is one of the widely-grown forest trees in South Asia [52]. Karanja seed contains 33 to 36% oil and 20 to 30% protein [53]. Nonedible energy crop ‘jatropha’, has received significant attention as one of the best candidates for biodiesel production due to its high oil content (50 to 60%) [50]. At the same time, jatropha seed kernel is rich in non-edible protein (27 to 32%) while kernel residue after oil extraction (seed cake) has a significantly high protein content (~60%) [44, 54]. Protein contents in the non-edible seed cakes are comparable or higher than that of commercial SF (~52%) [46]. Sunflower oil cake obtained from the oil industry is an inexpensive source of proteins which accounts for about 35% of protein in it [13]. It contains two major groups of protein, 2S albumin and 11S globulin and a significant amount of phenolic compounds. Castor bean cakes are also rich in protein. However, they cannot be used in food sector because of their toxicity. Recent research on these proteins has compared them with commercially available soy protein and found quite similar amino acid profiles [44]. The mechanical and thermal properties of various non-edible protein-based resins are presented in Table 2.2.

3–98

15–250

15–80

10–350

NA

14–59

Jatropha [46]

Karanja [41]

Neem [45]

Sunflower [48,55,56]

Castor bean [47]

Canola [31]

192–200

1.0–1.5

18–2000

54.9–705.1

226–711

460–2120

Young’s modulus (MPa)

* The properties show a broad range because of the modifiers used

Molecular weight (kDa)

Protein

Table 2.2 Properties of various non-edible protein sheets*.

8.2–11.9

3.6–6.2

2.8–12.5

3.6–13.3

7.8–17.6

8.5–22.1

Tensile strength (MPa)

17.5–30.0

28–96

2–140

22.9–103.4

5–39

1–26.3

Fracture strain (%)

50–65

NA

NA

86–103

NA

NA

Tg ( °C)

84–102

NA

NA

206–262

259

245.4

Td ( °C)

20 Advanced Green Composites

Green Resins and Strengthening Mechanisms 21

2.2.2 Plant Starch-Based Resins Plant-based starch is one of the most abundant naturally-occurring polymers on earth. It is inexpensive, biodegradable, environment-friendly, and yearly renewable. Commercial starches have been extracted from seeds, roots, and crops using simple separation processes. Thus far, both edible and non-edible starches have been utilized as resins in composite materials [7, 57–61]. Native starch consists of a mixture of two main macromolecular components: amylose and amylopectin, both made of D-glucose units. Amylose is a linear molecule obtained by α-D-(1 4) glycosidic bonds and 200–20,000 glucose units with a molecular weight of 105–106 that has excellent sheet forming ability whereas amylopectin is a highly branched molecule obtained by α-D-(1 4) linked by α-D-(1 6) glycosidic bonds in every 25–30 glucose units with a molecular weight of 107–109 [62–64]. Figure 2.4 shows chemical structures of amylose and amylopectin. Typically, native starches contain 15–30% amylose and 70–85% amylopectin [62, 63]. For example, most widely used starch-based resin for green composites are commercially-available corn and potato starches which contain 20–30% amylose with 70–80% amylopectin [62, 63]. A commonly used amylopectin-rich starch is waxy maize starch which contains around 99% amylopectin. However, there are some high-amylose starches available as well. These include amylomaize which contain as high as 50–80% of amylose [62]. The functional properties of native starches depend on their origin as well as the relative contents of amylose and amylopectin. The native starches exist in the form of granules which have dimensions ranging from 0.5 to 175 micrometers [62]. Starch granules differ in shape, size, structure, and chemical composition depending on the origin of the starch, organization of amylose and amylopectin, crystallinity, and molecular weight OH CH2OH

CH2OH

CH2OH

O

O

O

OH

OH OH

O

OH

O

OH

O OH

OH 300–600

OH

O

O

HO OH O

HO O

HO HO

O HO

O HO

Amylose

OH O HO

O HO

Amylopectin

Figure 2.4 Chemical structures of amylose and amylopectin.

O

22

Advanced Green Composites

[65]. Both amylose and amylopectin are structured in alternating layers of crystalline and non-crystalline regions by hydrogen bonding within the granules [66, 67]. As a result, most native starches exist in semi-crystalline form with crystallinity ranging between 15 and 45% [62]. Starch granules are insoluble in cold water. However, when immersed in hot water they absorb water, swell and, as a result, their crystalline structure gets disrupted. Water molecules then further interact with amylose and amylopectin, resulting in physical changes through further swelling and rupture of granules and partial solubilization of starch [68]. Heating starch solutions in excess solvents that have the ability to form hydrogen bonds at temperatures of 65–100 °C (depending on the type of starch) induces an irreversible process known as starch gelatinization. This process introduces irreversible changes in starch granules through two steps: the hydration and diffusion of the solvent into the starch granules, rupturing them and then melting the starch crystals [69]. Gelatinization process is highly affected by the kind of solvent and the starch/solvent ratio. Gelatinization makes the starch solutions more viscous and homogeneous which is advantageous for producing green resins for composite materials and other applications such as adhesives. Starch-based resin sheets have been obtained by two main methods: solvent evaporation or wet method and thermoplastic processing or dry method [63]. In the wet method, the starch powder is dispersed in a large amount of water along with various plasticizers such as polyols, sorbitol, and sucrose. Many authors have reported sheets obtained from a variety of starches using this method [58–60, 70–73]. The wet method involves four simple steps to prepare starch-based resin sheets: gelatinization, homogenization (in case of emulsions or mixtures), casting, and drying. The most essential condition to obtain starch-based sheets is that the granules of native starches must be ruptured or de-structured by the gelatinization process in the presence of excess water which is the first stage of the wet method. The gelatinization process is dependent on the origin of the starch due to the variation in their granule structures [74]. Hence, starch gelatinization takes place at different temperatures depending on starch, plasticizer, and water content, in most cases [63]. Although gelatinization occurs during heating of starch and water mixture, similar result can be obtained using an alkaline medium which is called cold gelatinization [63]. Once the gelatinization is accomplished, other compounds are mixed with it and this process is termed as ‘homogenization’. This step depends on the components in the sheet-forming solution. In the case of starch and plasticizers only, homogenization step may not be required. However, starch, with lipids or other immiscible components, must be processed through the

Green Resins and Strengthening Mechanisms 23 homogenization step to obtain a stable emulsion [75, 76]. The solutions are cooled and cast on leveled Teflon®-coated surfaces after gelatinization and homogenization and allowed to dry in controlled conditions to evaporate the excess water from the sheet. The wet method for starch sheet processing is not suitable at the industrial scale due to the long drying time. Dry method, however, is an alternative for producing starch-based sheets on a large industrial scale. In the dry method, starch-based sheets are obtained by using a thermal process in which the water content is much lower compared to the wet method. However, to be processed using dry method, the raw materials should have thermoplastic properties. Hence, the raw materials should become soft (melted or rubbery) at a temperature lower than their decomposition temperature and can be molded into a predetermined shape by a thermal/mechanical process. Although native starch is not a thermoplastic polymer, thermoplastic starch (TPS) can be produced by processing starch–plasticizer(s) mixtures in an extruder at temperatures of 140–160 °C using high pressure and repeated shear forces [77]. The presence of plasticizers with water is necessary in order to reduce the film/ sheet brittleness. Although the gelatinization process depends on water content and temperature, gelatinization of starches can also be obtained using low water content and high shear forces and pressure. High shear forces rupture the starch granules which results in faster water transfer into the starch molecules. Various dry methods for starch sheet processing have been discussed in detail as well as illustrated elsewhere. These include injection-molding [78], extruding with a film-blowing die [79], and thermo-pressing [80]. Starch sheets have low permeability to oxygen. However, they exhibit several drawbacks, such as hydrophilic character and poor mechanical properties. The major limitation of starch-based resin sheets is that they are very hydrophilic. As they absorb water, their properties deteriorate. Various chemical methods have been adopted to overcome the hydrophilicity problem such as modification of native starches through esterification, etherification, cross-linking, grafting, and enzyme treatment. A detailed review of the improvements of starch-based resin sheets can be found elsewhere [57]. Among these processes, etherification is the most common method to improve functional properties of resin on a commercial basis due to its low cost. Blending is another way to enhance the mechanical and barrier properties of starches. Blending of starch with biodegradable materials such as chitosan, agar, and PVOH have shown improved mechanical properties as reported earlier [32]. The functional properties of various starch-based resins are presented in Table 2.3. The properties show a very broad range due to a large variety of processing parameters and modifiers

26–28

71–73

39

20–1300

3–17

3.3–13

43–47

248–1592

50–53

Amylose (%)

Amylopectin (%)

Crystallinity (%)

Young’s modulus (MPa)

Tensile strength (MPa)

Fracture strain (%)

Glass transition ( °C)

O2 permeability (g.m-1s-1Pa-1)

WVP (g.m-1s-1Pa-1)** 6.3–13

6.4

152

2.1–85.2

3.5–35.5

3.9–2701

25

74–79

20–25

Potato

2.6–8.6

NA

NA

3–59

1.6–22.3

21–1052

35

74–86

14–26

Rice

NA

NA

-98–27

3.2–141

0.8–17.6

12–1354

39

99

1

Waxy

NA

NA

143

2.2–109

2.1–17.4

12–982

36

72–73

26–27

Wheat

Source: [59, 63, 72, 81–88]. *Several reports have shown the variation of the measured values; hence, a range has been quoted instead of a unique value. **WVP is the water vapor permeability which determines material’s ability to allow water vapor to pass through.

Corn

Starch

Table 2.3 Properties of various agricultural starch-based resin sheets*.

3.8–7.8

4.3–112

62–131

4.5–263

2.2–52.8

1–482

38

81–83

16–20

Cassava

24 Advanced Green Composites

Green Resins and Strengthening Mechanisms 25 used. However, the broad range of properties indicates possibilities for a broad range of applications. Similar to non-edible proteins, there have been a few studies focused on exploring non-edible starch sources as resin for composite materials since ethical questions have been raised on the utilization of edible starches for technical applications. Starch from some non-edible seeds, such as mango (Mangifera indica) has shown promising characteristics as an alternative to edible sources in green composites [60]. Mango seed cake, obtained after oil extraction, is a very rich source of carbohydrate (about 77%), most of which is starch that can be easily extracted [60]. Non-edible giant taro (Alocasia macrorrhiza), a tropical tuber crop, is also a good source of starch which may amount to over 85% of total dry matter [89]. Using such nonedible starches can be helpful in protecting the edible starches which are also used as a food source.

2.3 2.3.1

Green Resins from Microbial Fermentation Polyhydroxyalkanoates

PHAs and their copolymers are fully biodegradable and environmentfriendly aliphatic polyesters produced by microorganisms from renewable resources such as sugars and lipids. These bio-polyesters exhibit thermoplastic and elastomeric properties. PHAs, excellent alternatives to petroleum-derived resins such as PP and PE, have already been used successfully as matrices with bio-based fibers to produce fully ‘green’ composites. For example, the copolymer of PHAs, i.e., poly-b-hydroxybutyrate (PHB) and poly-b-hydroxybutyrate-co-valerate (PHBV) have been successfully used as matrices in various fiber-reinforced composites [90–93]. PHAs are synthesized by a wide variety of Gram-positive and Gram-negative bacteria cultivated using agricultural raw materials. PHAs are divided into two groups; short chain length and medium chain length PHAs, based on the number of carbon atoms in the chain of the branching polymers and the type of homopolymers or heteropolymers producing monomeric units. Short-chain length PHAs have 3–5 carbon atoms such as poly(3hydroxybutyrate) P(3HB), poly(4-hydroxybutyrate) P(4HB) and poly(3hydroxyvalerate) P(3HV) and the copolymer P(3HB-co-3HV) while medium-chain length PHAs consist of 6–14 carbon atoms or more than 14 carbon atoms such as poly (3-hydroxyhexanoate) P(3HHx), poly(3hydroxyoctanoate) P(3HO) and copolymers such as P(3HHx-co-3HO) [94]. Figure 2.5 shows the chemical structures of PH3B, PHV, and PHBV

26

Advanced Green Composites CH3 O

HC

O CH2C

PH3B

O

CH2CH3 O

HC

O

CH3

n

HC

CH2C

O PHBV

CH2CH3 O

CH

O

CH2C n

CH2C n

PHV

Figure 2.5 Chemical structures of PH3B, PHV, and PHBV.

[95]. The carbon chain length of the repeat units in PHAs including molecular weight varies depending on the bacteria, conditions of growth and method of extraction. The molecular weight can vary from about 50,000 to over a million. These thermoplastic polyesters of several R-hydroxyalkanoic acids are accumulated in intracellular granules by numerous microorganisms (Gram-negative and Gram-positive bacteria) in the presence of excess carbon, especially when another essential nutrient such as oxygen, nitrogen or phosphorus is limited [94]. These polyesters are produced as carbon and energy reserves or reducing power storage materials by the microorganisms. When a limiting nutrient is provided to the cell, these energy storage compounds are metabolized and used for bacterial growth as carbon source. PHAs differ in their physico-chemical properties depending on their chemical composition. Some of the general characteristics of PHAs are that they are water insoluble and relatively resistant to hydrolytic and UV degradation but has poor resistance to acids and bases. They are soluble in chloroform and other chlorinated hydrocarbons. Their hydrophobicity, melting point, glass transition temperature, and crystallinity depend on the monomer composition, although they are mostly hydrophobic. Hence, PHAs exhibit a wide variety of mechanical properties; from hard and stiff to elastic. For example, short-chain-length PHAs tend to be stiff, brittle and highly crystalline (around 60–80%) whereas medium chain-length PHAs have significant flexibility, elasticity with low crystallinity (about 25%), tensile strength, modulus, high elongation to break and low melting and glass transition temperatures [94].

Green Resins and Strengthening Mechanisms 27 The first discovered PHA in bacteria was PHB by Lemoigne in 1925 [95]. PHB is a linear polyester of D (-)-3-hydroxybutyric acid and is the most widely-utilized member of the PHA family. P(3HB) is highly crystalline biopolymer and is insoluble in water and, hence, resistant to hydrolytic degradation. Like starch, it also has low Oxygen (O2) permeability. However, it shows poor mechanical properties compared to petroleumbased polymers such as polypropylene. The density of PHB is in the range of 1.18–1.26 g/cm3 depending on the morphology. Typically, the molecular weight of P (3HB) is in the range of 10–3,000 kDa with a poly-dispersity index of 2. P(3HB) is very brittle and stiff material because of its high crystallinity. However, the physical properties of P(3HB) homopolymer could be improved by producing higher molecular weight polymers. Poly (4-hydroxybutyrate) P(4HB) is a strong thermoplastic material. The tensile properties of P(4HB) are comparable to those of PE. Also, it has significant elastic properties such as 100% fracture strain. However, the material properties of P(3HB) and P (4HB) can be manipulated by blending or combining with other polymers, or plasticizers such as glycerol, salicylic ester, soybean oil, epoxidized soybean oil, PE glycol, triacetine, fatty alcohol, dioctyl sebacate, which have been described in detail in an earlier review [95]. Another strategy to manipulate the functional properties of PHBs is to form copolymers by incorporating different HA monomers other than 3HB such as 3-hydroxyvalerate, 3-hydroxyhexanoate, 3-hydroxypropionate and 4-hydroxybutyrate into the polymer chain. Functional properties of P(3HB) such as crystallinity, melting point, stiffness and toughness have been improved significantly. One of the most successful copolymers is P(3HB-co-3HV) which has a lower crystallinity and melting temperature, increased elongation to break, lower modulus, and improved toughness compared to P(3HB) and similar to PP. Another copolymer Poly (3-hydroxybutyrate-co-3-hydroxyhexanoate) combines the functional properties of PE such as strength, ductility, and toughness with the physico-chemical properties of polyesters such as printability and dyeability. The average properties of PHAs are summarized in Table 2.4.

2.3.2 Pullulan Pullulan is a linear and water-soluble microbial polysaccharide based resin produced extracellularly using a variety of sugar-containing feedstocks by the black yeast-like fungus named Aureobasidium pullulans. It comprises of maltotriose units connected by α-1,4 glycosidic bond while consecutive maltotriose units are connected to each other by α-1,6 glycosidic linkages [98]. Figure 2.6 shows the chemical formula of pullulan. The molecular

28

Advanced Green Composites

Table 2.4 Properties of various PHA resins [94–97]. Properties

P(3HB)

P(4HB)

P(3HB-co- P(3HB-co- P(3HB-co4HV) 4HB) 4HA)

Density (g/cm3)

1.17–1.25

1.21

1.23–1.25

Glass transition (°C)

4–15

50

( 6)-6

42

Crystallinity (%)

60–70

NA

50–60

NA

NA

Young’s modulus (GPa)

3.5–4

0.15

0.7–2.9

0.03–0.1

0.2

Tensile strength (MPa)

15–40

104

30–38

24–65

17

Fracture strain (%)

3–8

1000

20

242–1080

680

Melting temp. (°C)

160–180

53

130–170

50–166

133

NA

NA 8

O CH2

CH2OH

CH2OH O

O OH

O

OH

OH

O

OH OH

O OH

O OH CH 2

Figure 2.6 Chemical formula of pullulan.

weight of pullulan varies from 45 to 600 kDa. The cultivation parameters such as temperature, pH, and types of carbon and nitrogen sources affect the molecular weight and, hence, the functional properties of the pullulan as a film. Pullulan has excellent film-forming ability and the film is biodegradable, non-toxic, colorless, tasteless, odorless, flexible, transparent, heat sealable, moldable, and impermeable to oil and oxygen [99]. Pullulan

Green Resins and Strengthening Mechanisms 29 decomposes at high temperatures in the range of 250–280 °C. It is highly soluble in water and alkali, however, insoluble in alcohol and some other organic solvents. Also, stable pullulan solution has relatively low viscosity compared to other natural polysaccharides. However, pure pullulan films have poor water resistance and mechanical performance. Although pullulan has several advantages as a matrix or resin in composite material, the high cost has limited the use of pullulan. One effective way to reduce the cost as well as to improve the functional properties of pullulan sheets is to blend it with inexpensive polymers that are compatible with it and have higher mechanical properties. Recent research has shown that blending pullulan with alginate, chitosan, cellulose, and starch improved its functional properties including thermal and mechanical properties, water vapor permeability and water absorption [100–103]. Also, composite films of pullulan and proteins, such as whey protein [104] and rice protein concentrate [105] have shown improved functional properties. Pullulangelatin films have demonstrated higher tensile strength combined with decreased oxygen permeability and lower cost [106], while the use of rice wax has shown improvements in water vapor barrier properties [107].

2.4 Green Resins Using Monomers from Agricultural Resources 2.4.1

Polylactic Acid

PLA is a thermoplastic biopolymer produced from lactic acid (2-hydroxy propionic acid) as the basic building block. Figure  2.7 shows the chemical structures of PLA. Lactic acid (2-hydroxy propionic acid) is prepared through bacterial fermentation process of agricultural products such as corn, rice, potatoes, sugar beet, etc. It is the simplest hydroxy acid with an asymmetric carbon atom and exists in two optically active configurations. The L-stereoisomer (L-lactic acid) is produced in humans and other mammals while both D-enantiomers are produced in bacterial systems. Similar to PHAs, it is an aliphatic polyester-based resin and it can be semi-crystalline or fully amorphous. PLA has been utilized extensively as a green resin with bio-based fibers to produce fully green composites [108–112]. It has become a very popular bio-resin as an alternative to petroleum-derived polymers due to its comparable functional properties. Detailed reviews have been reported on the properties of PLA-based green composites elsewhere [113, 114]. PLA is classified into two groups; low molecular weight and high molecular weight PLA [115]. Low molecular-weight PLA can be synthesized by

30

Advanced Green Composites O

O

O

O

HO

HO OH H

H3 C

CH3

(a)

H

(b)

C

H C

OH

n

CH3

(c)

Figure 2.7 Chemical structures of (a) L-lactic acid, (b) D-lactic acid ad (c) Polylactic acid (PLA).

CH3 HO n

OH

Dehy dra

tive c

Lactic acid

O

Lower molecular weight prepolymer MW 2,000 to 10,000

n de

n

HO

CH3

O

tio

Co

OH

O

n

sa

CH3

CH3

O O

onde

–H O

O

2

Chain coupling agents

CH3

nsatio

n

HO –H2O

CH3

O O n

Condensation

CH3

O

OH

O O

High molecular weight PLA MW –100,000

O

CH3

CH3

O

HO

n

O

OH

O

CH3

Lower molecular weight prepolymer MW 1,000 to 5,000

Ring openeing polymerization

O Deplo y

meriz

ation

O

H3C

O

CH3

O

O

Lactide

Figure 2.8 Typical polymerization processes for Polylactic acid (PLA) [115].

direct condensation of lactic acid. However, high-molecular weight PLA is prepared either by condensation/coupling of lactic acid, azeotropic condensation of lactic acid or ring opening polymerization of the cyclic lactide dimer. Typical polymerization processes are presented in Figure 2.8 [115]. In condensation polymerization, eliminating water as condensation product is critical for the equilibrium reaction to proceed in the forward direction. At first, the free water present in the agricultural products is separated by evaporation and simple sugars from agricultural products (for example; glucose, maltose, and dextrose from corn or potatoes and sucrose from cane or beet sugar) are extracted. In the second step, lactic acid is obtained

Green Resins and Strengthening Mechanisms 31 through fermentation of these simple sugars. The self-condensation of lactic acid results in a low-molecular-weight product. Although the condensation polymerization is the least expensive route, it is difficult to obtain high molecular weights in the solvent-free medium. As a result, the use of coupling agents is required. However, that adds cost as well as complexity. The azeotropic condensation polymerization is another method to obtain high molecular-weight PLA without the use of coupling agents. In this method, the polycondensation reaction is performed in an organic solvent medium. The water generated during the process along with the organic solvents is separated using molecular sieves or drying agents. The polymer is either isolated or dissolved and precipitated for further purification. A third method, ring-opening polymerization, is industrially more preferred method to obtain high molecular weight PLA. In this method, lactic acid obtained through fermentation of the simple sugars is converted into lactide in the presence of a catalyst. After purification by vacuum distillation, lactide is converted into PLA polymer through the ring opening polymerization in the presence of a suitable catalyst. PLA is a fully biodegradable resin requiring a degradation time in the range of six months to two years in the soil burial. After composting, PLA-based materials are converted into water and carbon dioxide which are consumed in growing more agricultural products for further conversion to PLA. PLA has good mechanical properties that are comparable to polyethylene terephthalate (PET) and PP which are the most common polymers used in automobile applications. For example, semicrystalline PLA has a tensile modulus of 3 GPa along with a tensile strength of 50 to 70 MPa. In addition, the flexural properties of PLA are also suitable as resin for the structural composites. The melting temperature of PLA (near 180 °C) along with the available standard processing equipment is safe for combining it with natural fibers since natural fibers do not degrade at these processing temperatures. Like most thermoplastic polymers, PLA can be processed by extrusion, injection-molding and sheet casting, blow molding and electro-spinning which make it a versatile material. PLA is a hydrophobic polymer due to the presence of the –CH3 side chain. Because of these favorable properties, PLA has been a strong candidate among the biopolymers as a green resin to be used in green composites [113, 114]. The mechanical, thermal and permeability properties of PLA resin are listed in Table 2.5. However, using hydrophilic natural cellulosic fibers as reinforcement has been a little difficult since the fiber/PLA bonding is weak. To improve fiber/resin bonding and to obtain good mechanical properties of the composites either coupling agents are used or fibers are treated to improve their surface chemistry and topography.

32

Advanced Green Composites

Table 2.5 Properties of PLA resin [116]. PLA Density (g/cm3)

1.24–1.29

Glass transition (°C)

55–80

Crystallinity (%)

37

Young’s modulus (GPa)

2.7–3.8

Tensile strength (MPa)

30–70

Fracture strain (%)

4–7

O2 permeability (g.m 1s 1Pa 1)

4.38 × 10–18

WVP (g.m 1s 1Pa 1)

1.34 × 10–14

Melting temp. (°C)

173–178

2.5 Strengthening of Green Resins using Nano-Fillers Although the advantages of green resins such as their biodegradability, environment-friendliness, and renewability have made them highly attractive for use as an alternative to the currently used petroleum-based resins, low mechanical properties of some of the resins still limit their utilization in various load-bearing structural applications. In the past two decades, however, researchers have successfully enhanced various green resin properties by incorporating nano-fillers such as nanoclay, calcium carbonate, hydroxyapatite (HA), micro-fibrillated cellulose, chitin nanowhiskers, cellulose nanowhiskers, microcrystalline cellulose (MCC), and carbon nanotubes, etc. [24–26, 117–121]. It has been shown that nano-fillers have high potential to improve the mechanical, physicochemical as well as thermal properties of biopolymers [24, 25, 46]. They exhibit a high specific surface area, as well as high strength and stiffness, which make them potential candidates for reinforcing biopolymers [122]. The effects of nano-fillers on the functional properties of polymers depend on their size, shape, surface characteristics and, importantly, the degree of dispersion. Higher specific surface area of nano-fillers as compared to microparticles provides significantly higher interface for stress transfer and lower stress concentration factors in the composite. To get the full potential of nano-fillers, the key is to be able to transfer the potential functional properties of the nano-fillers to the polymer. In this context,

Green Resins and Strengthening Mechanisms 33 two main issues need to be solved. These issues are; interfacial bonding and uniform dispersion of the individual nano-fillers in the polymer. However, due to their high specific surface area and energy, nano-fillers induce undesirable van der Waals and electrostatic forces between them which can result in large agglomerations [123]. These large agglomerations produce unwanted stress concentration factors which act as failure or crack initiators in nanocomposites and eventually, deteriorate the mechanical properties of the composites [124]. As a result, uniform dispersion is a prime requirement to overcome the van der Waals attractions between the nano-fillers. Another significant advantage of the uniform dispersion of individual nano-fillers is that they increase the tortuosity of the resin molecules, increasing the resin modulus. Various mechanical methods to disperse nano-fillers in polymer resins, such as high speed stirring, ultrasonication and high shear mixing have been reported in the literature [124]. Chemical functionalization of nanofiller surfaces has also improved their dispersion into the resin [125]. Even when the dispersion issue is overcome, poor interfacial compatibility can cause the nano-fillers to separate from the polymeric resin during processing and/or in use. In order to fully exploit the potential of nano-fillers as structural reinforcement of polymer, a strong nanoparticle/resin interfacial adhesion is desired. The interfacial adhesion between nano-fillers and resin can be improved by functionalizing the surface chemically [126]. A wide variety of nano-fillers have been utilized as promising reinforcements to enhance the mechanical performance of green resins. Typically, the nano-fillers used in green resins to improve the mechanical properties can be divided into two categories: inorganic nano-fillers and organic nano-fillers. In the next sections, some bio-based inorganic and organic nano-fillers that are commonly used in fully green composites, without compromising their environmental advantages, have been discussed.

2.5.1 Inorganic Nano-Fillers Layered silicates or clays are one of the most important inorganic nanofillers that have received enormous interest as reinforcement for polymeric resin due to their wide availability and low cost. Generally, clay-based nanofillers are also referred to as aluminosilicate-based nano-fillers. Among these nano-fillers, montmorillonite is the most utilized reinforcement in green composites since they are environment-friendly, economical, and naturally abundant [127]. Also, they have superior mechanical properties and chemical resistance [80, 119, 128]. To some lower extent, halloysite nanotubes (HNT) as fibrillar silicate have been interesting nanofiller

34

Advanced Green Composites AI, Fe, Mg, Li OH Tetrahedral

O Li, Na Rb, Cs

Octahedral

Tetrahedral

Exchangeable cations

Figure 2.9 The general structure of clay-based nano-fillers [127].

in green nanocomposites [129–131]. The general structure of clay-based nano-fillers is comprised of sheets arranged in structural layers in which each layer is formed by 2, 3 or 4 sheets. The sheets are formed either by tetrahedral (T) [SiO4]4 or by octahedral (O) [AlO3(OH)3]6 arrangement. Figure 2.9 shows the general structure of clay-based nano-fillers [127]. All clay-based nano-fillers are composed either of TO or of TOT layers while each layer has a thickness of about 1 nm and lateral dimensions in the range between 300 Å and several micrometers. The size and shape of nano-fillers depend on the origin, source and method of preparation. Structurally, the clay-reinforced polymeric nanocomposites can be intercalated or exfoliated nanocomposites. Figure  2.10 shows the schematic illustration of two different types of thermodynamically achievable clayreinforced nanocomposites. In the case of intercalated nanocomposites, polymer chains intercalate (or place) themselves between individual layers by expanding while layers are completely delaminated in the polymer in exfoliated nanocomposites. As mentioned earlier, one key issue is the dispersion of clays into the polymer. There are various ways to obtain uniform dispersion in the green resins. However, the dispersion and distribution of clays in the polymer

Green Resins and Strengthening Mechanisms 35

Intercalated

Exfoliated

Figure 2.10 Schematic illustration of two different types of thermodynamically achievable clay-reinforced nanocomposites [128].

has been achieved more effectively through the chemical modification of nano-fillers and manipulation of the processing parameters. Researchers have successfully dispersed clay-based nano-fillers in various green resins such as PLA, soy protein, and starch to improve the mechanical properties of these resins; particularly their stiffness and strength [119, 128, 132–134]. A list of nano-fillers which have successfully improved the properties of many green resins after their incorporation is provided in Table 2.6. Another natural and most-abundantly available nanofiller is calcium carbonate (CaCO3). It has been investigated as reinforcement for many polymers in the past decade [135]. Researchers have been successful in tailoring the properties of some green resins; particularly their stiffness and toughness, by incorporating CaCO3 nanoparticles to fabricate advanced inorganic particulate-filled polymeric nanocomposites [24, 118, 136, 137]. More successful attempts have been with soy protein and PLA based resins where the resins modified using CaCO3 nanoparticles have shown significant improvements in their mechanical properties [24, 136–139]. HA or calcium phosphate [Ca10(PO4)6(OH)2] nanoparticles have also received attention as nano-fillers to enhance mechanical and thermal properties of green resins [140–142]. Particularly, it has been a suitable reinforcing constituent in composite materials for biomedical applications such as tissue scaffolds and bone regenerations because of its excellent functional properties including biocompatibility, osteoconductivity, non-toxicity and bioactivity [140, 143]. HA, in fact, is the prime inorganic constituent (65– 70%) of bone and hard tissue matrix that is indirectly bound to collagen protein through some non-collagenous proteins such as osteocalcin, osteoponcin or osteonectin in nanocrystal form [122]. It has two binding sites, Ca2+ and PO43 which possess affinity towards biological macromolecules [144]. An alternative economical and environment-friendly approach to synthesizing CaCO3 and HA nanoparticles has been developed from

SPC

MMT nanoclay

PHB

Potato starch

Potato starch

Cassava starch

Corn starch

Corn starch

SPI

Resin

Nano-fillers (NF)

3060 3440

5

195.6

5 0

29.8

376

5 0

188

–

2.5 0

–

825

5 0

790

206.74

30 0

38.15

587.6

20 0

180.2

2124

7 0

717

Young’s modulus (GPa)

0

NF (%)

27.0

29.6

5.2

3.3

9.82

7.33

108.18

13.26

–

–

27.34

5.51

14.5

8.8

74.5

50.1

Tensile strength (MPa)

Table 2.6 Mechanical properties of green resins reinforced with various inorganic nano-fillers.

–

–

46.8

62.6

44

68

1.88

1.79

12

10

17.82

85.32

5

90

9.5

14.8

Fracture strain (%)

[149]

[148]

[147]

[146]

[134]

[133]

[132]

[119]

References

36 Advanced Green Composites

CaCO3

PLA

Halloysite nanotubes

PLA

SPI

Potato starch

Wheat starch

SPI

Hydroxyapatite (HA)

PHBV4

PHB

PHB

PHB-HV

2190 3890

10

1055

5 0

675

700

9 0

100

31.3

8 0

12.8

3644

6 0

2923

1737

5 0

923

1971

4.4 0

1689

2201

5 0

1884

4250

4 0

4130

795

3 0

481

0

63.0

51.5

28.3

22.6

43

27

2.93

2.28

70.3

54.5

41

25

29

23

28

24

28.6

32.9

33

31

5.5

4.0

11.2

24.8

6

30

83

87

3.2

3.6

7.2

18

1.97

1.6

1.47

1.56

0.74

0.9

5.6

8.5

[139]

[24]

[131]

[130]

[129]

[153]

[152]

[152]

[151]

[150]

Green Resins and Strengthening Mechanisms 37

Advanced Green Composites

38

various bio-wastes such as egg shells, sea shells, bovine bones, fish bones, etc., since these wastes are rich sources of calcium precursors [24, 25]. It has been reported that more than 45 million kg of chicken eggshell waste which contains more than 94% CaCO3 is disposed of every year by the food industry in the United States without further processing, creating environmental pollution [145]. The unique chemical composition and almost unlimited availability of the eggshell has made it a viable and inexpensive (almost free) source for obtaining bio-based CaCO3 to synthesize HA nanoparticles while simultaneously reducing the eggshell waste.

2.5.2 Organic Nano-Fillers Cellulose, discovered and isolated by Payen (1838), is the most abundant naturally-occurring macromolecule on earth with an estimated production of 1.5 1012 tons/yr [154]. It is a linear homopolymer chain of β-Dglucopyranose units linked by β-(1 4) glycosidic bonds. The repeat unit of cellulose consists of 2 glucose (monomer) units and is known as cellobiose (C12H22O11). The number of cellobiose units in cellulose molecules is in the range of 5,000 to 7,500, making the degree of polymerization (DP) of 10,000 to 15,000, depending on the source or origin. Figure 2.11 shows the basic structural unit of cellulose (cellobiose) and the structure of cellulose strand. Purity, DP and crystallinity also vary among the cellulose varieties depending on the source. Most cellulose fibers or wood varieties contain cellulose nanocrystals (CNC) or nanofibrils within which the molecules are highly oriented. As a result, nanocrystals and nanofibrils have high H O

OH O HO

O HO

OH HO O

O

OH OH

O

H O

n

O HO H O O HO H O

(a)

(b)

O HO

O OH O OH O OH O OH

HO O O H HO O O H HO O O H HO O O H

OH

H O

O

O HO

OH O

H O O HO

OH

H O

O

O HO

OH

H O

O

O HO

O OH O OH O OH O OH

HO O O H HO O O H HO O O H HO O O H

OH O

OH O

OH O

OH O

Figure 2.11 (a) Basic structure of cellulose and (b) structure of a strand of cellulose.

Green Resins and Strengthening Mechanisms 39 strength and stiffness. Because of the high mechanical properties combined with low density, biodegradability and renewability, nano-scale cellulose production and their application in the fabrication of strong ‘green’ composites have received increasing attention [6]. Microfibrillated cellulose (MFC), introduced in early 1980s by Turbak et al. [155] and Herrick et al. [156], has been a promising nanofiller for green resins to improve their mechanical properties. MFC can be obtained from many varieties of cellulose materials through a two-step mechanical shearing process, namely initial refining step followed by a high-pressure homogenization. In general, the homogenization process where cellulose fiber slurry is passed through a small orifice under a high pressure produces high shear forces that cause transverse cleavage or separation along the longitudinal axis of the cellulose microfibrillar structure. During each pass, the material gets sheared further to become smaller and more uniform in diameter, resulting in fibrils with high aspect ratio. Obtained MFC is typically in the range of 10 to 500 nm in diameter and usually consists of aggregates or entangled cellulose nanofibrils (~3–10 nm) [157]. The aspect ratio of MFC can be high, 100 to 150 or higher, since the length varies in the range of 0.5–10 μm and the diameter varies in the range of 10–100 nm [6]. MFC is very pure cellulose (~100%) and while it has both amorphous and crystalline regions, it has high crystallinity. Due to its high specific strength, modulus, and aspect ratio, MFC has been used as a reinforcing agent for a broad range of biopolymers including proteins and ‘green’ composites [6, 154, 158–160]. However, MFC tends to flocculate through hydrogen bonding due to its strong hydrophilic character derived from the hydroxyl groups which results in a non-uniform dispersion even in plant-based hydrophilic resins, thus, resulting in lower than the expected ultimate mechanical properties of ‘green’ nanocomposites [6, 154]. Further, its high aspect ratio makes it easy to get entangled. Full potential of MFC as reinforcement in ‘green’ nanocomposites, however, can only be achieved by homogeneous and unidirectional dispersion of MFC in the resin. MCC is cellulose-based nano or microfiller that has been used to reinforce green resins such as PLA and SPI [161,162]. It is produced by acid hydrolysis of cellulose such as wood pulp, the least expensive sustainable resource [163]. During the hydrolysis process, the amorphous part of cellulose is hydrolyzed first, dissolved and removed. This results in tightly packed microcrystals to be separated. In the next step, mechanical method is used to disintegrate the microcrystals. These disintegrated microcrystals aggregate again after drying due to strong hydrogen bonds and form porous and spongy MCC nano-fillers.

40

Advanced Green Composites

Cellulose nanowhiskers or nanocrystalline cellulose are rodlike CNC. Their width varies between 5 and 70 nm and the length can be in the range of 100 nm to several micrometers [154]. CNC is generated by acid hydrolysis of cellulosic biomass such as wood pulp, sugar beet, MCC, cotton, tunicate, wheat straw, etc., in mineral acids such as sulfuric or hydrochloric acid [154]. At first, cellulose pulps are milled to uniform small particles and hydrolyzed to remove lignin, hemicellulose, and the amorphous cellulose part which separates the crystalline fibrils. Acid hydrolysis is then terminated by rapid dilution and the acid is removed by centrifugation and/or dialysis. The disintegration of cellulose fibrils is then achieved by ultra-sonication process to obtain a suspension of CNC. CNC is highly crystalline material with a degree of crystallinity of 90% or greater [163]. The dimensions of CNC depend on the morphology of the raw material and the process used. The width is observed in the range of 3–10 nm while the length varies from 100 nm to several micrometers [163]. The reported elastic modulus is in the range of 143–150 GPa for single CNC crystals [163]. High aspect ratio as well as high mechanical properties of CNC makes them valuable for use in enhancing the properties of green resins. Significant efforts have been reported on the use of CNC, MCC and MFC to reinforce nanocomposites and several ways of obtaining uniform dispersion of cellulose nanofibrils in various green resins to obtain reinforced nanoresins [46, 162, 164–167]. Table  2.7 presents the mechanical properties of various green resins reinforced with CNC, MCC and MFC. The data in Table  2.7 clearly indicate that adding cellulosic nano-fillers enhance the mechanical properties significantly. A major difficulty in cellulose-reinforced composites, however, is the compatibility of reinforcement and resins to achieve the full potential of nanoscale cellulose. As mentioned earlier, uniform dispersion of the reinforcement in the resin is critical to obtain the highest possible properties. Without compatibility with the resin, cellulose reinforcements tend to aggregate, reducing their effectiveness. In addition, good resin/reinforcement interfacial interaction or bonding is needed to achieve substantially higher mechanical properties. Nano-scale cellulose has the same noncompatibility issue as MFC but at a greater scale due to its high specific surface area and high surface energy which create agglomerates easily. The other factors to improve the mechanical properties of the resins or composites are the aspect ratio of nanoscale cellulose and the composite preparation method. Since cellulose is hydrophilic, it is difficult to achieve good dispersion in hydrophobic resins such as PLA and PHAs. However, several methods have been reported for dispersing these nano-fillers in green

Waxy maize

MFC/NFC

PLA

PLA

PLA

Cassava starch

SPC

Mango starch

Jatropha protein

Resin

Nano-fillers (NF)

2900

9000

70 0

5000

5700

20 10

4000

84.3

5 0

44.5

2374

40 0

589

2407

40 0

1347

1147

20 0

720

1686

15 0

1354

Young’s modulus (GPa)

0

Content NF (%)

58.9

105

35

69.4

60.9

4.8

4.1

74.9

21.7

50.03

16.24

20.3

15.1

44

17.6

Tensile strength (MPa)

Table 2.7 Mechanical properties of green resins reinforced with various organic nano-fillers.

3.4

2.5

1

1.7

3.1

71

83.3

10.2

14.4

2.65

1.54

6.3

8.2

7.8

3.1

Fracture strain (%)

(Continued)

[170]

[164]

[165]

[169]

[168]

[61]

[46]

[59]

References

Green Resins and Strengthening Mechanisms 41

MCC

NCC

Nano-fillers (NF)

Table 2.7 Cont.

SPI-Gelatin

PLA

PLA

PLA

PHBV

Potato starch

Resin

47.52 120.03

3.5

3110

5 0

2650

5000

25 0

3600

3710

5 0

3310

1760

5 0

820

6200

70 0

1.6

3600

5 0

Young’s modulus (GPa)

Content NF (%)

5.94

1.70

52.4

62.8

36.2

49.6

66.5

56.1

26.1

14.1

160

0.35

71.2

Tensile strength (MPa)

25.7

95.2

3.1

19.5

1.6

2.4

2.3

2

7.8

12.4

8.1

80

2.7

Fracture strain (%)

[162]

[172]

[161]

[171]

[167]

[157]

References

42 Advanced Green Composites

Green Resins and Strengthening Mechanisms 43 resins among which physical methods, such as calendering and thermal and non-thermal methods such as corona, plasma, and laser treatments are successfully employed [163]. However, chemical modification methods which modify the cellulose surfaces by surface compatibilization by grafting desired chemical groups and/or co-polymerization, have been the most widely studied and preferred methods [163].

2.6 Conclusions The development of eco-friendly and biodegradable polymers and nanocomposites is best viewed in the context of greening the technology for ecological and environmental benefits. These green materials are designed to degrade when disposed of in normal environmental conditions such as composting, by the actions of living organisms and thus do not have to end up in landfills, like petroleum-derived materials do. Besides, research efforts in the development of these polymers have shown promising opportunities towards replacing conventional petroleum-derived polymers in various nonstructural applications such as casing, packaging, furniture, automotive panels, and other indoor secondary structures in housing and transportation, etc. Still, there are substantial opportunities to achieve further improvements in their functional properties and thus broadening the applications, particularly in primary structures where fibers such as Kevlar®, glass and graphite are commonly used as reinforcement. Also, global commercialization of green materials to replace conventional petroleum-based ones is a challenge since the conventional technology is mature and the products have been developed over several decades. In addition, they are inexpensive and familiar to the customer. Hence, new and novel pathways need to be developed for obtaining strong and tough eco-friendly green polymers and composites inexpensively and increase their usage while reducing the overall material carbon footprint at the same time. In addition, non-edible resin sources can be explored from a range of forestry wastes such as defatted castor, linseed, palm seed cakes and many others which do not have much value. Various vegetable/food and fishery wastes can also be viable sources for protein, starch or cellulose based resins for future ‘green’ composites. Sources for synthesizing inorganic nanomaterials such as HA, calcium carbonate, etc., such as animal bones, egg shells and teeth from slaughterhouses as well as from dead animals can also be explored on commercial scale. Thus, bio-derived resins and nanomaterials from waste sources can be manipulated for the reduction of edible source utilization and improvement of waste management.

44

Advanced Green Composites

Non-edible seedcakes are good sources of sugar-based polysaccharides as well. These sugars can be modified for direct use as plasticizers or modified to sugar derivatives such as multifunctional aldehydes or acids through oxidation which can be good candidates as ‘green’ cross-linkers. Thus, both the resins and cross-linkers can be derived and/or synthesized from the same source. Development of ‘green’ materials is still in its infancy, but the progress in the past decade has been fast. In the future, the consumption of the green and sustainably derived materials can only be expected to increase as the technology matures and becomes affordable.

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153. M. M. Rahman, A. N. Netravali, B. J. Tiimob, V. Apalangya, and V. K. Rangari, Bio-inspired “green” nanocomposite using hydroxyapatite synthesized from eggshell waste and soy protein, J. Appl. Polym. Sci. 133, 2016. 154. D. Klemm, B. Heublein, H. P. Fink, and A. Bohn, Cellulose: Fascinating biopolymer and sustainable raw material, Angew. Chem. Int. Ed. 44, 3358–3393, 2005. 155. A. F. Turbak, F. W. Snyder, and K. R. Sandberg, Microfibrillated cellulose, a new cellulose product: Properties, uses, and commercial potential, J. Appl. Polym. Sci. Appl Polym Symp. 37, 815–827, 1983. 156. F. W. Herrick, R. L. Casebier, J. K. Hamilton, and K. R. Sandberg, Microfibrillated cellulose: Morphology and accessibility, J. Appl. Polym. Sci. Appl. Polym. Symp. 37, 797–813, 1983. 157. A. J. Svagan, M. A. S. Azizi Samir, and L. A. Berglund, Biomimetic polysaccharide nanocomposites of high cellulose content and high toughness, Biomacromolecules 8, 2556–2563, 2007. 158. I. Siró and D. Plackett, Microfibrillated cellulose and new nanocomposite materials: A review, Cellulose 17, 459–494, 2010. 159. H. P. S. A. Khalil, A. H. Bhat, and A. F. I. Yusra, Green composites from sustainable cellulose nanofibrils: A review, Carbohydr. Polym. 87, 963–979, 2012. 160. G. Siqueira, J. Bras, and A. Dufresne, Cellulosic bionanocomposites: A review of preparation, properties and applications, Polymers 2, 728–765, 2010. 161. A. P. Mathew, K. Oksman, and M. Sain, Mechanical properties of biodegradable composites from poly lactic acid (PLA) and microcrystalline cellulose (MCC), J. Appl. Polym. Sci. 97, 2014–2025, 2005. 162. C. Li, J. Luo, Z. Qin, H. Chen, Q. Gao, and J. Li, Mechanical and thermal properties of microcrystalline cellulose-reinforced soy protein isolate– gelatin eco-friendly films, RSC Adv. 5, 56518–56525, 2015. 163. C. Miao and W. Y. Hamad, Cellulose reinforced polymer composites and nanocomposites: A critical review, Cellulose 20, 2221–2262, 2013. 164. A. N. Nakagaito, A. Fujimura, T. Sakai, Y. Hama, and H. Yano, Production of microfibrillated cellulose (MFC)-reinforced polylactic acid (PLA) nanocomposites from sheets obtained by a papermaking-like process, Compos. Sci. Technol. 69, 1293–1297, 2009. 165. L. Suryanegara, A. N. Nakagaito, and H. Yano, The effect of crystallization of PLA on the thermal and mechanical properties of microfibrillated cellulosereinforced PLA composites, Compos. Sci. Technol. 69, 1187–1192, 2009. 166. A. Iwatake, M. Nogi, and H. Yano, Cellulose nanofiber-reinforced polylactic acid, Compos. Sci. Technol. 68, 2103–2106, 2008. 167. L. Jiang, E. Morelius, J. Zhang, M. Wolcott, and J. Holbery, Study of the Poly(3-hydroxybutyrate-co-3-hydroxyvalerate)/Cellulose nanowhisker composites prepared by solution casting and melt processing, J. Compos. Mater. 42, 2629–2645, 2008

Green Resins and Strengthening Mechanisms 55 168. X. Huang and A. N. Netravali, Environmentally friendly green materials from plant-based resources: Modification of soy protein using gellan and micro/nano-fibrillated cellulose, J. Macromol. Sci. Part A. 45, 899–906, 2008. 169. E. de M. Teixeira, D. Pasquini, A.A.S. Curvelo, E. Corradini, M.N. Belgacem, and A. Dufresne, Cassava bagasse cellulose nanofibrils reinforced thermoplastic cassava starch, Carbohydr. Polym. 78, 422–431, 2009. 170. M. Jonoobi, J. Harun, A. P. Mathew, and K. Oksman, Mechanical properties of cellulose nanofiber (CNF) reinforced polylactic acid (PLA) prepared by twin screw extrusion, Compos. Sci. Technol. 70, 1742–1747, 2010. 171. D. Bondeson and K. Oksman, Polylactic acid/cellulose whisker nanocomposites modified by polyvinyl alcohol, Compos. Part Appl. Sci. Manuf. 38, 2486–2492, 2007. 172. D. Bondeson and K. Oksman, Dispersion and characteristics of surfactant modified cellulose whiskers nanocomposites, Compos. Interfaces 14, 617–630, 2007.

3 High Strength Cellulosic Fibers from Liquid Crystalline Solutions Yuxiang Huang1 and Jonathan Y. Chen2* 1

2

Chinese Academy of Forestry, Beijing, China The University of Texas at Austin, Austin, TX, USA

Abstract Viscose rayon process used to manufacture cellulosic fiber for over a century is uncompetitive due to its lengthy procedure, high water and air pollution, and high-energy consumption. Since it was reported that concentrated solutions of cellulose derivative could self-assemble into ordered phases, research in the area of liquid crystal formation of cellulose using new solvent systems has been primarily driven to develop new nonpolluting processes to produce regenerated cellulosic fibers. This chapter provides a concise overview on high strength cellulosic fibers that can be regenerated from liquid crystalline solutions of cellulose derivatives and nonderivatized cellulose, as well as new methods to reinforce the cellulose liquid crystalline solutions. Keywords: Cellulose, regenerated cellulose, liquid crystalline solution, anisotropic solution, ionic liquid solvent, solution spinning, cellulose/nanoparticle composite

3.1 Introduction As the most plentiful and renewable polymer provided by nature, cellulose can be obtained from wood, cotton, ramie, hemp, flax, jute, kenaf, bagasse, sorghum, switch grass, and innumerable other sources of lignocellulose biomass. At the turn of the last century, the production of cellulose was approximately 100 billion metric tons per year [1]. Due to its low cost,

*Corresponding author: [email protected] Anil Netravali (ed.) Advanced Green Composites, (57–66) © 2018 Scrivener Publishing LLC

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environmental friendliness, biodegradability, nontoxicity and biocompatibility, cellulose is widely used in food, clothing, housing, paper production, packing, pharmaceutical industries, and many other fields. Most recently, much attention has been drawn to biodegradable cellulosic fibers as reinforcement in ‘green’ composite materials that can be used in wide ranging applications such as transportation, aerospace, sporting goods, civil engineering structures and many others [2–5]. However, polymer composites reinforced with plant-based cellulosic fibers have relatively moderate mechanical properties such as stiffness and strength on account of the relatively low tensile properties of plant-based cellulose fibers which range between 300 and 900 MPa for tensile strength and 10 to 30 GPa for Young’s modulus (stiffness) [6]. Plant-based fibers also suffer from wide variation in their properties which can depend on the region they are produced, the variety, temperature, rain fall fluctuations, and other factors. In addition to these, the plant-based fibers are short in length. Spinning them into continuous forms such as yarns, by twisting, reduces their strength and stiffness even further. Therefore, it is much desired to produce continuous cellulose fibers with inherently higher stiffness, strength and evenness to improve the mechanical properties of ‘green’ composites and, if possible, to make them ‘advanced green composites’. It is known that liquid crystalline solutions of polymers are good precursors for preparing high tenacity/modulus fibers. Fibers produced from liquid crystalline solutions, e.g., aramids (Kevlar® being the prime example) have outstanding mechanical properties due to their intrinsic polymer structure, high crystal orientation, rigid backbones and strong intermolecular secondary bonds [7]. In fiber spinning from solutions of flexible polymers or isotropic melts, an after-spinning treatment, such as drawing, is always required to improve molecular orientation and, thus, their tensile properties. However, this post-spinning treatment is no longer a necessity when using liquid crystalline solutions, because the local orientation order of the molecular chains already exists at such high level that they can be easily transformed into highly oriented fibers [8]. This is because the spinning processes used such as air-gap wet spinning is capable of retaining the molecular/crystalline orientation in the spun fibers. Initially, polymers with stiff backbone were used for preparing liquid crystalline solutions, most of which consisted of aromatic units such as poly (p-phenylene terephthalamide) or PpPTA, polybenzoxazole or PBO, poly-benzothiazole or PBT, and poly{2,6-di-imidazo[4,5-b:4 5 -e]pyridinylene-1,4-(2,5-dihydroxy)phenylene} or PIPD. Cellulose-based polymers have also been used to produce liquid crystalline solutions depending on whether the polymers were from derivatized cellulose or from non-derivatized cellulose.

High Strength Cellulosic Fibers 59

3.2 Fibers from Liquid Crystalline Solutions of Cellulose Derivatives Cellulose derivatives which have a better solubility than cellulose can form liquid crystals when dissolved in suitable solvents [9]. The cellulose bone or polymer is a single-stranded or linear homopolymer with β-linked 1,4-anhydroglucose units. Each anhydroglucose unit contains three hydroxyl groups which provide convenient sites for substitution reactions which can lead to a wide variety of cellulose ethers and esters. Figure 3.1 shows an example of a heavily substituted cellulose chain. Many of these derivatives form cholesteric liquid-crystalline phases in suitable solvents, and some derivatives with relatively large but non-mesogenic side-chains form both lyotropic and thermotropic liquid crystals [10]. Originally, Pannar and Willcox from DuPont filed a patent application on lyotropic systems of various cellulose derivatives in a great variety of solvents [10]. After that, Werbowyj and Gray dealt with hydroxypropylcellulose in water, which was the first publication on a cellulose-based liquid crystalline system [11]. Since then many more mesophases of solutions of cellulose derivatives in a number of solvents have been described. These include cellulose ester and ethers in both organic solvents and inorganic acids, cellulose acetate in trifluoroacetic acid, ethylcellulose in acrylic acid, etc. The conditions for mesophase formation of cellulose-based systems OH

H3C C H2C

H O

H2C

O

O

CH2 C

H3C

O

O O

H

CH3 O C O CH3 C H H H2C O O

OH

O CH2 O O

C

O

CH3

H C CH3

O

CH2 CH2 O

H C H3 C

O

C CH3

Figure 3.1 Illustration of a cellulose derivative: (acetoxypropyl) cellulose composed of anhydroglucose units with ether and ester substituents.1

1

Reproduced from Faraday Discuss. Chem. Soc., 79, D. G. Gray, Chemical characteristics of cellulosic liquid crystals, 257–264, 1985, with permission from The Royal Society of Chemistry.

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can be summarized as follows [10]. Use of strongly acidic or polar solvents can promote cellulose mesophase formation. Highly substituted cellulose in which the substituent is a large group, forms nematic phases in many solvents while cellulose with fewer or smaller substituents needs specific solvent interaction to form mesophases. When the liquid crystalline solutions were prepared, they were usually spun through solution (wet) spinning with an air gap, in which stretching was applied, into an aqueous coagulation bath [8, 13, 14]. Northolt et al. developed high modulus/high strength cellulose fibers that were spun from an anisotropic solution in phosphoric acid [8]. In their process, powdered cellulose was dissolved in a mixed solution of orthophosphoric and polyphosphoric acid in which the phosphorous pentoxide (P2O5) concentration approximated to 74% w/w. The solution was then filtered, heated and extruded through an assembly of spinnerets for fiber spinning. The formed filaments passed through an air gap, in which they were subjected to stretching, and went through a falling jet coagulation bath of acetone for cellulose regeneration. After washing the remaining acids, the produced filaments were neutralized with Na2CO3 solution. The regenerated cellulose fibers produced using this approach featured high crystallinity as well as high crystalline orientation. As a result of that, these fibers had a tensile strength in the range of 1.7 GPa, highest cellulose fiber strength at the time, and tensile modulus in the range of 58 GPa. Such strong fibers have been found to be suitable for advanced green composites with high strength and toughness [13, 14]. It has also been found that chemical and mechanical treatments can further increase their orientation and crystallinity and in turn, their tensile properties [14].

3.3 Fibers from Liquid Crystalline Solution of Nonderivatized Cellulose As mentioned above, due to the poor solubility of cellulose, it needs special solvents. Many reports have shown successful use of N-methylmorpholineN-oxide (NMMO)/water, trifluoroacetic acid/dichloromethane, and phosphoric acid, etc., to obtain anisotropic solutions of cellulose [8, 15, 16]. Recently, ionic liquids (ILs), a group of new organic salts that exist as liquids at a relatively low temperature, have been used as novel solvents for dissolving cellulose [12, 17, 18]. They have many attractive properties, such as chemical and thermal stability, non-flammability and immeasurably low vapor pressure. They are considered as ‘green’ solvents and have been widely used compared with traditional volatile organic compounds that can pollute the environment. Some studies have shown that cellulose

High Strength Cellulosic Fibers 61 without derivatization can be well dissolved in some hydrophilic ILs, such as 1-butyl-3-methylimidazolium chloride (BMIMCl), 1-allyl-3-methylimidazolium chloride (AMIMCl) and 1-ethyl-3-methylimidazolium (EMIMAc). Among them, EMIMAc has superior solubility for cellulose than the other two ILs [9]. It has been found that cellulose/ILs solutions appear optically anisotropic when cellulose dissolves at high concentrations [18, 19]. Based on this, Song et al. found that the two critical cellulose concentrations for the appearance of biphase and fully anisotropic phase for microcrystalline cellulose (MCC)/AMIMCl solutions were 9 and 16 wt% [7]. For MCC/EMIMAc solutions, lyotropic liquid crystalline phase formed at 10 wt% cellulose concentration. Fiber spinning of cellulose using ILs offers a major advantage of a simple, one-step dissolution process without the use of aggressive solvents as compared to the traditional viscose rayon process. Jiang et al. dissolved cotton pulp using BMIMCl by heating at 90 °C and then extruded the solution through a spinneret at various spinning speeds [20]. The results showed that spinning speed had a significant effect on fiber properties. Increasing spinning speeds increased both crystallinity and crystal orientation, thereby enhancing the fiber tenacity (tensile strength) and tensile (Young’s) modulus. Kosan et al. used several ILs including NMMO-MH, BMIMCl, EMIMCl, AMIMCl, BMIMAc and EMIMAc to dissolve Eucalyptus pre-hydrolysis sulfate pulp and cotton linter pulp and spin fibers using those solutions [21]. Regenerated cellulose fibers produced from the IL solvents had similar tensile strength but lower elongation and modulus compared to the properties of regenerated cellulose fibers produced by the NMMO system. It was also found that cellulose fibers produced using acetate anions had a lower tenacity but higher elongation than those produced using chloride anions [21]. It is also important to know whether the structure of regenerated cellulose polymer is the same or different from that of raw cellulose. Sun et al. reported their recent study on this issue using wide angle X-ray diffraction (WAXD) technique [22]. They found that raw cellulose from a southern pine pulp showed typical Cellulose I structure (Figure 3.2). After BMIMCl dissolution and fiber spinning, the regenerated cellulose revealed a dominant Cellulose II profile (Figure 3.3).

3.4 Regenerated-Cellulose/CNT Composite Fibers with Ionic Liquids Most recently, it was recognized that the addition of nano-fillers greatly affected the phase structure of a polymer solution, such as liquid crystalline

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Intensity (A.U.)

Experimental Model-total Model-amorphous Othters: model-peaks

22.48 (2 0 0)

14.91 (1 1 0) 16.56 (1 1 0)

0

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20

30

40

50

60

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Figure 3.2 Raw pulp cellulose structure by WAXD pattern.2

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19.84 (1 1 0)

11.98 ( 1 0)

0

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30 40 2 (degree)

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Figure 3.3 Regenerated cellulose structure by WAXD pattern.3

2

3

Reprinted from Carbohydrate Polymers, 118, Sun, L., Chen, J.Y., Jiang, W., and Lynch, V., Crystalline Characteristics of Cellulose Fiber and Film Regenerated from Ionic Liquid Solution, 150–155, 2015, with permission from Elsevier. Reprinted from Carbohydrate Polymers, 118, Sun, L., Chen, J.Y., Jiang, W., and Lynch, V., Crystalline Characteristics of Cellulose Fiber and Film Regenerated from Ionic Liquid Solution, 150–155, 2015, with permission from Elsevier.

High Strength Cellulosic Fibers 63 phase, which can be formed when anisotropic nanoparticles were combined with an isotropic polymer solutions [23]. As a result, the anisotropic nanoparticles (such as carbon nanotubes, graphene oxide) have been used as key potential precursors for the fluid phase processing of particles into aligned materials with outstanding properties [24]. Luo et al. found that 9 wt% of MCC/AMIMCl solution resulted in an isotropic solution [25]. They further found that upon adding 0.63 wt% halloysite nanotubes (HNTs), the solution resulted in an ordered liquid crystalline structure. The tensile strength of the regenerated fibers spun from MCC/AMIMCl solution with the addition of 7 wt% HNTs was 130.1 MPa, not as strong as those fibers obtained from liquid crystalline cellulose solutions obtained in phosphoric acid [8]. Rahatekar et al. used multiwall carbon nanotube (MWCNT)/Cellulose/ EMIAc solution to spin cellulose based nanocomposites [26]. The cellulose/ MWCNT composite fibers had a tensile strength of 257 MPa at 0.07 mass fraction of MWCNTs, still significantly lower than 1.7 GPa obtained for the fibers obtained from liquid crystalline cellulose solutions obtained in phosphoric acid [8]. Zhang et al. also obtained reinforced regenerated cellulose fibers with a tensile strength of 335 MPa by adding 4 wt% MWCNTs [27]. All of these studies found that higher loadings of CNTs in composite fibers did not perform as well as lower loadings due to the increase in voids generated or other defects as the loading of CNTs increased [26]. Furthermore, the addition of nanoparticles into cellulose solutions ultimately affects cellulose polymer structure during the cellulose regeneration. Chen et al. have reported on this phenomenon in their recently published paper [28]. For example, when producing cellulose/nanosilver fiber, BMIMCl solvent was used with nanosilver loading between 0.5 and 2% mass fraction. Their results indicated that the addition of Ag nanoparticles into the cellulose solution could increase the fiber crystallinity but decrease the fiber crystallite size significantly. This is because when the cellulose solution was extruded into an aqueous bath, cellulose was regenerated as water (antisolvent) entered into the solution system to reduce solvent dissolvability. The cellulose regeneration began with cellulose nucleation and growth. The presence of nanosilver increased the number of crystal nucleation seeds and created larger number of crystallites but reduced the time allowing for crystallite growth.

3.5 Future Prospects Impacted by the development of bioenergy and bio-economy, comprehensive utilization of lignocellulosic biomass for green material production

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has become more appealing and competitive. Synthetic polymeric fibers derived from petroleum feedstock are facing a challenge in terms of their sustainability, though they largely dominate the end-use markets for textiles, apparel, consumer products, and engineering materials at present. Liquid crystalline solutions of plant cellulose have provided green material manufacturers with a new platform for producing high strength cellulosic fibers. Especially, the dissolution of cellulose using ILs provides the possibilities of preparation of various advanced composites, including cellulose derivatives and cellulose composites. Although research in this area is still not mature, once these fibers are fully developed they will have a bright future. Advanced green composites with excellent strength and stiffness (above 1000 MPa and over 37 GPa, respectively) have already been fabricated using liquid crystalline cellulose fibers [25, 26], which can be selected for use in primary structural applications. It is expected that the high strength cellulosic fibers from liquid crystalline solutions used for advanced green composites will experience a new level of development in the future.

Summary As the people’s will for natural resource conservation and environment protection has helped shape up a biobased economy featuring sustainable manufacture, preservative consumption, and eco-friendly urban development, biobased lignocellulosic fiber has gained a new attraction as an alternative to petroleum based non-renewable fibers that are used in many applications today. However, most plant-based cellulose fibers are weak in strength and their Young’s modulus is low compared to synthetic fibers. This hinders the application of cellulose fibers in the high performance green composites. Therefore, cellulose fibers with inherently higher strength and stiffness are critical to overcome the mechanical weakness of the green composites. This chapter provides a concise overview of the high strength cellulosic fibers that can be prepared from liquid crystalline solutions of cellulose derivatives and non-derivatized cellulose. Additionally, nanoparticles can be added into cellulose solutions to reinforce the regenerated cellulose fibers. Cellulose fibers produced from such liquid crystalline solutions have excellent tensile properties and can be expected to promote the future development of advanced green composites with high strength and stiffness and in some cases high toughness as well.

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References 1. M. Lima and R. Borsali, Rodlike cellulose microcrystals: Structure, properties, and applications, Macromol. Rapid. Comm. 25, 771–787, 2004. 2. M. L. Hassan, E. A. Hassan, and K. N. Oksman, Effect of pretreatment of bagasse fibers on the properties of chitosan/microfibrillated cellulose nanocomposites, J. Mater. Sci. 46, 173–140, 2011. 3. W. Liu, M. Misra, P. Askeland, L. T. Drzal, and A. K. Mohanty, ‘Green’ composites from soy based plastic and pineapple leaf fiber: Fabrication and properties evaluation, Polymer 46, 2710–2721, 2005. 4. M. J. John, B. Francis, K. T. Varughese, and S. Thomas, Effect of chemical modification on properties of hybrid fiber biocomposites, Commpos. Pt. A-Appl. Sci. Manuf. 39, 352–363, 2008. 5. M. J. John and S. Thomas, Biofibres and biocomposites, Carbohyd. Polym. 71, 343–364, 2008. 6. A. K. Mohanty, M. Misra, and G. Hinrichsen, Biofibres, biodegradable polymers and biocomposites: An overview, Macromol. Mater. Eng. 276–277, 1–24, 2000. 7. H. Song, J. Zhang, Y. Niu, and Z. Wang, Phase transition and rhelogical behaviors of concentrated cellulose/ionic liquid, J. Phys. Chem. B 114, 6006–6013, 2010. 8. M. G. Northolt, H. Boerstoel, H. Maatman, R. Huisman, J. Veurink, and H. Elzerman, The structure and properties of cellulose fibres spun from an anisotropic phosphoric acid solution, Polymer 42, 8249–8264, 2001. 9. H. Song, Y. Niu, Z. Wang, and J. Zhang, Liquid crystalline phase and gelsol transitions for concentrated microcrystalline cellulose (MCC)/1-ethyl3-methylimidazolium acetate (EMIMAc) solutions, Biomacromolecules 12, 1087–1096, 2011. 10. D. G. Gray. Chemical characteristics of cellulosic liquid crystals, Faraday Discuss. Chem. Soc. 79, 257–264, 1985. 11. P. Manuel and W. O. Burr, Optisch anisotrope spinnmassen und verwendung derselben., German Patent DE2705382, 1977. 12. R. S. Werbowyj and D. G. Gray, Liquid crystalline structure in aqueous hydroxypropyl cellulose solutions, Mol. Cryst. Liq. Cryst. 34, 97–103, 1976. 13. S. Zhu, Y. Wu, Q. Chen, Z. Yu, C. Wang, S. Jin, Y. Ding, and G. Wu, Dissolution of cellulose with ionic liquids and its application: a mini-review, Green Chem. 37, 325–327, 2006. 14. J. T. Kim and A. N. Netravali, Fabrication of advanced “green” composites using potassium hydroxide (KOH) treated liquid crystalline (LC) cellulose fibers, J. Mater. Sci. 3950–3957, 2013. 15. A. N. Netravali, X. Huang, and K. Mizuta, Advanced ‘green’ composites, Adv. Composite Mater. 16, 269–282, 2007.

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16. D. L. Patel and R. D. Gilbert, Lyotropic mesomorphic formation of cellulose in trifluoroacetic acid-chlorinated-alkane solvent mixtures at room temperature, J. Polym. Sci. Polym. Phys. Ed. 19, 1231–1236, 1981. 17. H. Chanzy, A. Peguy, S. Chaunis and P. Monzie, Oriented cellulose films and fibers from a mesophase system, J. Polym. Sci. Polym. Phys. Ed. 18, 1137–1144, 1980. 18. N. V. Plechkova and K. R. Seddon, Applications of ionic liquids in the chemical industry, Chem. Soc. Rev. 37, 123–150, 2008. 19. H. Zhang, J. Wu, J. Zhang, and J. He, 1-Allyl-3-methylimidazolium chloride room temperature ionic liquid: A new and powerful nonderivatizing solvent for cellulose, Macromolecules 38, 8272–8277, 2005. 20. R. P. Swatloski, S. K. Spear, J. D. Holbrey, and R. D. Rogers, Dissolution of cellose with ionic liquids, J. Am. Chem. Soc. 124, 4974–4975, 2002. 21. G. Jiang, Y. Yuan, B. Wang, X. Yin, K. S. Mukuze, W. Huang, Y. Zhang, and H. Wang, Analysis of regenerated cellulose fibers with ionic liquids as a solvent as spinning speed is increased, Cellulose 19, 1075–1083, 2012. 22. B. Kosan, C. Michels, and F. Meister, Dissolution and forming of cellulose with ionic liquids, Cellulose 15, 59–66, 2008. 23. L. Sun, J. Y. Chen, W. Jiang, and V. M. Lynch, Crystalline characteristics of cellulose fiber and film regenerated from ionic liquid solution, Carbohyd. Polym. 118, 150–155, 2015. 24. V. V. Makarova, M. Y. Tolstykh, S. J. Picken, E. Mendes, and V. G. Kulichikhin, Rheology-structure interrelationships of hydroxypropylcellulose liquid crystal solutions and their nanocomposites under flow, Macromolecules 46, 1144– 1157, 2013. 25. C. M. Koo, S. O. Kim, and I. J. Chung, Study on morphology evolution, orientational behavior, and anisotropic phase formation of highly filled polymerlayered silicate nanocomposites, Macromolecules 36, 2748–2757, 2003. 26. Z. Luo, A. Wang, C. Wang, W. Qin, N. Zhao, H. Song, and J. Gao, Liquid crystalline phase behavior and fiber spinning of cellulose/ionic liquid/halloysite nanotubes dispersions, J. Mater. Chem. A 2, 7327–7336, 2014. 27. S. S. Rahatekar, A. Rasheed, R. Jain, M. Zammarano, K. K. Koziol, A. H. Windle, J. W. Gilman, and S. Kumar, Solution spinning of cellulose carbon nanotube composites using room temperature ionic liquids, Polymer 50, 4577–4583, 2009. 28. H. Zhang, Z. G. Wang, Z. N. Zhang, J. Wu, J. Zhang, and J. S. He, Regenerated-cellulose/multiwalled-carbon-nanotube composite fibers with enhanced mechanical properties prepared with the ionic liquid 1-Allyl-3methylimidazolium Chloride, Adv. Mater. 19, 698–704, 2007. 29. J. Y. Chen, L. Sun, W. Jiang, and V. M. Lynch, Antimicrobial regenerated cellulose/nano-silver fiber without leaching, J. Bioact. Compat. Pol. 30, 17–33, 2015.

4 Cellulose Nanofibers: Electrospinning and Nanocellulose Self-Assemblies You-Lo Hsieh Department of Fiber and Polymer Science, University of California, Davis, USA

Abstract This chapter highlights top-down and bottom-up approaches to generate ultra-fine cellulose fibers of nano-scale dimensions, micro-porous, meso-porous and sheathcore hybrid structures as well as surface functionalized fibrous materials. Versatile solvent systems have been established to efficiently dissolve cellulose acetate to enable robust electrospinning into homogeneous 1D nanofibers and submicron size fibers that could be easily converted from amorphous to moderately crystalline cellulose II via alkaline hydrolysis. By pairing with either compatible or incompatible polymers, hybrid and nanocomposite fibers as well as porous fibers may be engineered. Surface reactions, grafting and electrostatic deposition can further offer surface functionalization for controlled hydrophobicity, stimuli-responsive behaviors and bound enzyme catalyses. Highly crystalline I nanocelluloses with varied geometries and surface chemistries have been efficiently derived via chemical means and/or shear forces in nanorod and nanofibril forms that have been shown to exhibit unique dispersing and emulsifying properties for oil-water emulsions, coagulants for microbes as well as templates for nanoparticles and nanoprisms. These nanocelluloses can be facilely assembled into new fibrous structures, super-absorbent hydrogels, films and amphiphilic to hydrophobic aerogels for applications, such as oil recovery and separation, water purification, etc. While ultra-fine fibers from these two approaches share some common fibrous morphologies, their crystalline structures, thermal behaviors, surface chemistires, reactivity and chemical functionalities are distinctively different, offering a wide range of strategies for fabricating cellulose nanofibers with tunable functional characteristics for novel materials and advanced composites. Keywords: Electrospinning, nanocellulose, self-assembly, morphology, surface

Corresponding author: [email protected] Anil Netravali (ed.) Advanced Green Composites, (67–96) © 2018 Scrivener Publishing LLC

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4.1 Introduction Nanoscale materials are of great interest for their novel structural properties and their high specific surface, i.e., surface area per volume or mass, among others. Nanoparticles or zero dimensional (0 D) nanomaterials of inorganic origins have been most widely studied for their ease of synthesis and unique surface characteristics as nanometer scales approaching molecular scales. However, their particulate sizes also attribute to their easy dissemination into the environment and biological systems such as human and marine lives, raising serious concerns. One dimensional (1 D) nanomaterials, such as nanorods and, in particular, nanofibers with higher aspect ratios, can be more easily integrated into larger and well inter-connected structures, thus more easily handled or potentially less likely to migrate into the surrounding environment and biological systems. Strictly speaking, nanomaterials refer to those materials with dimensions less than 100 nm. Nanofibers are 1 D nanomaterials with their lateral dimensions less than 100 nm. However, the term “nanofibers” have been commonly used for those with lateral dimensions, i.e., widths, thickness or diameters, in the hundreds nm. Nanofibers that are tens of nanometer wide are three orders of magnitude narrower than the conventional fibers with typically tens of micrometer widths; therefore their specific surfaces are three orders of magnitude higher. Such ultra-high specific surface fibrous materials are highly desirable for applications that rely on surface properties, such as controlled release and delivery, separation and filtration of chemicals, biological and pharmacologically active agents and molecules, etc. At these nanometer scale sizes, new properties and unexpected behaviors may also surface, offering potentially new functions for applications. Cellulose, nature’s most abundant polymer, is among the most renewable materials to be targeted for nanomaterials. Native cellulose found in a wide range of origins is highly fibrillated and semi-crystalline, which makes it difficult to dissolve and recalcitrant to be refined and converted into biofuels and small molecules. Cellulose is a long chain polysaccharide, poly(1,4 -anhydroglocopyranose), that consists of D(+)-glucopyranose building blocks linked by (1→4)-glucosidic bonds (Figure 4.1a). The -1,4 link between C1 and C4 of the adjacent D-glucose units resembles the stereo-regularity of a syndiotactic polymer, as opposed to the isotactic form of starch amylose with -1,4 link. The cellobiose structural repeating units provide the steric effects to limit free rotation of the C-O-C link, adding stiffness to the already bulky glucose structure. The three hydroxyl groups on each anhydroglucose unit, i.e., one primary C6 and two secondary C2 and C3, along with the chain

Cellulose Nanofibers

8.201 b 10.380

7.784a 8.201 96.5°

OH O O HO (a)

b

OH

OH

HO O

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O O OH (b)

(c)

Figure 4.1 Cellulose: (a) cellubiose repeating unit; cellulose I unit cell (b) top view along C axis; (c) along A axis [1] (Reprinted with permission from Industrial & Engineering Chemistry Research. Copyright 2015 American Chemical Society).

conformation allow optimal inter-molecular and intra-molecular hydrogen bonding to further enhance the rigidity of the cellulose chain and crystalline structure. In fact, the highly hydrogen bonded crystalline structure of cellulose is the main impediment to hydrolysis to fermentable sugars for biofuel generation. Native cellulose from higher plants consists of crystalline domains in cellulose Iβ form (Figure 4.1b) [1] (two chain monoclinic unit cell: a = 7.784 Å, b = 8.201 Å, c = 10.380 Å, α = β = 90°, γ = 96.5°) [2] whereas that from bacteria is in cellulose I form. While early effort in utilizing cellulose focused on dissolution of cellulose for regenerated cellulose fibers or cellulose derivatives for fibers and films, more recent interest has been on separating and using the crystalline domains or the so-called nanocelluloses. The interest in nanocelluloses has increased tremendously and especially over the past decade as evident by not only the increased numbers of papers but also reviews on the subject. The most studied source of cellulose to date is wood pulp whereas efforts related to other plant sources, especially agriculture crop residues and food processing by-products, are much less. This chapter presents some current perspectives and summarizes some of the recent developments related to nanocelluloses from non-wood sources without duplicating the reviews already available in the literature. Furthermore, the attention is drawn to the processes by which nanocelluloses are assembled into nanofibers. To date, nanofibers have been most commonly fabricated by electrospinning in which either polymer solutions or melts are ejected by high voltages, after overcoming their surface tensions, into unstable jets that further reduce in sizes to generate nanoscale fibers with extremely long lengths. Fabricating nanofibers by electrospinning has been attractive because of easy fiber formation from a wide range of organic polymers including those that cannot be spun into fibers by conventional solution

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or melt spinning processes due to lacking of specific viscoelastic properties and/or chain length requirements. However, electrospinning tends to produce fibers in heterogeneous widths that are, more often than not, in the hundreds of nm to a few micrometers. Electrospinning is generally lower in productivity than conventional fiber spinning processes, but is often off set by the unique advantages in fiber qualities it offers and/or by applications as a minor component with other bulk materials such as in the case of coating or laminating on other fibrous or porous substrates for filtration applications.

4.2 Electrospinning of Cellulose Solutions As cellulose does not melt and is difficult to dissolve into homogeneous solutions, electrospinning of cellulose is challenging and has been met with limited success. While wet spinning of cellulose has been commercialized for over a century and continuing research in cellulose dissolution has advanced and reviewed recently [3], wet processing remains to be the key approach. While cellulose can be dissolved in N-methyl-morpholine N-oxide (NMMO), lithium chloride/dimethyl acetamide (LiCl/DMAc) and ionic liquids, electrospinning cellulose from these solvents into fibers have been met with various dielectric and solvent removal constraints, among others. Electropsinning of cellulose from NMMO can produce 3–10 mm wide fibers [4] but requires elevated temperatures (70–110 °C) to lower viscosity [5] or smaller 200–700 nm wide fibers by using water as a coagulant to diffuse the solvent [6, 7]. While LiCl/DMAc dissolves cellulose well and charge carrying LiCl salt facilitates electrospinning into fibers, the volatile DMAc and the salt must be removed in subsequent steps [8, 9]. Ionic liquid, such as 1-butyl-3-methylimidazolium chloride, also dissolves and enables electrospinning of cellulose, but the solvent must be removed by coagulation in ethanol [10]. In all cases of electrospinning cellulose from these solvents, neither the feasibility of continuous fiber production nor the productivities of electrospinning from these solvent systems were reported.

4.3 Cellulose Nanofibers via Electrospinning and Hydrolysis of Cellulose Acetate A facile approach to cellulose dissolution and fiber spinning is to work with cellulose derivatives, in particular cellulose esters that are mass produced

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as commodity polymers and are commercially available. By replacing the pendant hydroxyls with ester groups, inter-molecular hydrogen bonding capability is essentially eliminated while interaction with organic solvents is significantly enabled. Significant lowering of chain lengths to 300–500 units of glucose repeating units or degree of polymerization that are about one order of magnitude lower than the cellulose precursor further enhance the solubility. In the case of cellulose acetates (CAs), solubility in a range of organic liquids becomes possible. CA with a molecular weight of 30 kDa and a degree of substitution (DS) of 2.45, i.e., an average of 2.45 of 3 hydroxyls being replaced with esters, is soluble in solvents with Hildebrand solubility parameters ( ) between 9.5 and 12.5 (cal/cm3)1/2 such as acetone, acetic acid and N,N-dimethylacetamide (DMAc) and dimethylforamide (DMF). However, electrospinning CA (30 kDa) in any of these liquids does not produce uniform fibers [11, 12]. Mixing DMAc with either acetone or acetic acid, both having lowered surface tensions and boiling points, enables efficient electrospinning of CA [9]. For instance, 15% CA in acetone/DMAc mixtures in mass ratios between 2:1 and 10:1 could be continuously electrospun into fibrous membranes. The 2:1 acetone/ DMAc mixture has shown to be the most versatile, permitting 12.5–20% CA solutions to be continuously electrospun into highly uniform fibers with smooth surfaces deposited in thousands of layers across the thickness of the fibrous mat (Figure 4.2) and in increasing widths from about 100 nm to 2 μm with increasing CA concentrations. These concentrations of CA in 2:1 acetone/DMAc correspond to η between 1.2 and 10.2 poise and around 26 dyne/cm. Solutions with viscosities lower than this lead to extensive bead formation whereas those with higher viscosities result in instability in electrospinning. The fiber packing may also be regulated by collecting fibers on different materials, e.g., more densely packed fibers on conductive aluminum foil or

50 m

(a)

(b)

Figure 4.2 Electrospun CA (DS=2.45, 30 kDa, 20 w%) in 2:1 acetone/DMAc: (a) top and planar view, (b) cross-sectional view.

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(b)

(c)

Figure 4.3 Electrospun cellulose acetate (CA, DS=2.45, 30 kDa) from 2:1 acetone/DMAc: (a) 20% CA collected on paper; (b) 20% CA collected on water; (c) 15 % CA collected on water, showing fiber packing as affected by collecting materials (a & b) and fiber size as affected by solution concentrations (b & c) [11] (Reprinted with permission from Journal of Polymer Science Part B: Polymer Physics. Copyright © 2002 Wiley Periodicals, Inc.).

water and less packed fibers on non-conductive cellulose paper (Figure 4.3). In all cases, these CA fibers may be deacetylated in mild (0.05 M) aq. or ethanolic NaOH, to DS 1.14 in 140 min or DS 0.15 in 60 min, respectively, with the latter more efficient and complete. This approach of productive electrospinning CA in versatile solvent mixtures and efficient hydrolysis is significant and has been followed by others in the following years [13–16]. Other cellulose derivatives, such as hydroxypropyl cellulose, hydroxypropyl methyl cellulose, and ethyl-cyanoethyl cellulose, can also be dissolved in solvents suitable for electrospinning but are not reverted back to cellulose to be included in this discussion. In addition, fibrous membranes for many novel functional properties and broad applications are discussed in the following section.

4.4 Bicomponent Hybrid and Porous Cellulose Nanofibers With similar physical constants, such as a high dielectric constant desirable for electrospinning as DMAc, DMF also dissolves CA. CAs of higher molecular weights of 50 and 60 kDa can be electrospun from DMF alone into uniform fibers whereas the lower 30 kDa CA requires the addition of dioxane to be electrospun into nanofibers in 1:1 DMF/dioxane mixture [17]. It should be noted that CA cannot be electrospun in dioxane alone either, similar to acetone, acetic acid and DMAc. In DMF, the semidilute entangled concentration (Ce) that is commonly used to determine the onset concentration for fiber spinning begins at 19.6% for 30 kDa CA and 16.5% for 50 kDa CA) CA. The fact that fiber formation from 50 kDa CA in

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DMF is more efficient at 20% concentration, significantly higher than the 16.5% Ce, suggests that other solution behaviors also contribute to electrospinning. Never-the-less, fibers electrospun from 20% CA (30 kDa) in 1:1 DMF/dioxane mixture were 50 to 200 nm wide, much smaller than those from the 2:1 acetone/DMAc mixture (0.83 to 2.5 mm). This is attributed to the low dielectric constant and higher compatibility of dioxane with CA than acetone. This significant reduction in fiber sizes demonstrates the critical role of solvents in not only regulating electrospinning but also the fiber structure. DMF is also a common solvent for many synthetic polymers, thus amenable to incorporate other polymers, such as poly(vinyl pyrrolidone) (PVP) [17], -cyclodextrin ( -CD) [17] and poly(ethylene oxide) (PEO) [18], with CA. In the cases of PVP and -CD, their additions greatly facilitate electrospinning of CA. The charge-holding ability of PVP enables electrospinning of CA/PVP mixtures in DMF alone into uniform hybrid nanofibers containing up to 70% CA (30 kDa) that by itself cannot be electrospun from DMF. In fact, PVP is not only electrospun into fibers from DMF alone, but its addition to CA increases the electrospinning rate of CA/PVP mixtures to approach that of PVP alone. CA/PVP bicomponent hybrid nanofibers consist of phase-separated PVP domains in a continuous CA matrix in increasing widths from 20 to 650 nm, with either higher PVP molecular weights (55 to 360 kDa) or quantities (up to 70 %) [15]. Upon dissolving PVP in water, mesoporous CA nanofibers with rough surfaces can be fabricated from CA/PVP hybrids with as high as up to 50% PVP. With 50 kDa CA, a 12.5 wt% solution in DMF is not as electrospinnable because it is below the minimum concentration necessary for chain entanglement (Ce), but can be efficiently electrospun into nanofibers when equal amount of -CD was added [17]. -CD is 7-membered truncated cyclic oligosaccharide with all hydroxyls pointing outwards from the truncated ring: C6 primary hydroxyls along the smaller rim and C2 and C3 secondary hydroxyls along the larger rim. The fact that -CD enables fiber formation of CA below its Ce concentration indicates that -CD molecules are well dispersed among CA chains as individual molecules to enhance CA chain entanglement via hydrogen bonding interactions between remaining hydroxyls of CA and -CD hydroxyls as confirmed by Fourier transform infrared (FTIR) spectroscopy. Up to 50 wt% of -CD can be well distributed in the CA matrix as individual inclusive complexes with 2-nm mesopores, exactly internal cavity dimension of -CD. As -CD has a hydrophobic cavity, non-polar compounds of similar or smaller dimensions may be encapsulated within cellulose nanofibers in significant proportion for robust controlled release applications.

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While both CA and PEO are easily soluble in DMF, the DMF-dioxane binary solvent system supports efficient electrospinning of all combinations of CA (30–50 kDa) and PEO (10–600 kDa) into CA/PEO bicomponent hybrid nanofibers [18]. Generally, a higher molecular weight of either CA or PEO lowers the threshold concentration necessary to achieve sufficient viscosities for fiber formation by electrospinning and results in generating wider fibers. Increasing total CA/PEO concentrations (5–20%) improves fiber formation by electrospinning, but also leads to wider fibers. In most cases, PEO phase separated into sheath with CA continuous matrix core structures. Another advantage of these versatile solvent systems is the ability to incorporate nanomaterials with CA. Multi-walled carbon nanotubes (MWCNTs) can be well dispersed with CA in 2:1 wt:wt acetone/DMAc to be electrospun, then alkaline hydrolyzed into MWCNTs-filled cellulose nanofibers [19]. Both the addition of MWCNTs and hydrolysis reduce fiber widths, by up to ca. 50 and 100 nm, respectively. While, not surprisingly, only about a third of the fibers appear to contain MWCNTs at very low loadings of 0.11 and 0.55 wt%, all MWCNTs are well aligned along the fiber direction, showing shear force during electrospinning to be sufficient to orient MWCNT nanofibers within CA matrix. In fact, the addition of 0.55 wt% MWCNTs also reduced CA fiber widths from 321 to 228 nm. All as-spun CA fibrous membranes, with or without MWCNT, are amorphous whereas the hydrolyzed cellulose nanofibers show two crystalline peaks, one at 20.2° peak, close to the typical 101 reflection of cellulose II (2θ = 19.8°) another at 22.0°, close to the 002 reflection of cellulose I (2θ = 22.5°) [19]. However, both are far weaker than the cellulose I observed in cotton [20] and cellulose II in regenerated cellulose [21]. These XRD patterns suggest that low quantities of cellulose I and II crystalline allomorphs are present in the fully hydrolyzed cellulose nanofibers and the presence of MWCNTs appears to slightly enhance their crystallinity. Intriguingly, both strength and wetting properties of cellulose nanofibers were significantly improved even at such low MWCNT loadings.

4.5 Wholly Polysaccharide Cellulose/Chitin/Chitosan Hybrid Nanofibers Chitin is the second most abundant polysaccharide in nature. Like cellulose, chitin and chitosan have been widely studied as platforms for advanced materials. Together, cellulose, chitin and chitosan have been combined to take the advantage of their collective desirable properties.

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Chitin has the same β-1,4-D(+)-anhydroglusidic main chain with two hydroxyls (C-6 primary and C-3 secondary) as cellulose but differs only on the C-2 pendant group, i.e., an acetyl amine as opposed to a hydroxyl for cellulose. Hydrolyzing the acetyl amine to amine converts chitin to chitosan, which is the most significant cationic polysaccharide with many desirable properties. Also, like cellulose, the rigid main chain structure and the extensive inter- and intra-molecular hydrogen bonded structure of chitin and chitosan give them their non-thermoplastic characteristics and insolubility in most common solvents, significant barriers to their processing and conversion to materials. Thus, processing chitin or chitosan into fibers is also challenging. For instance, chitosan is soluble in 1:1 acetone/acetic acid at up to 3%, but none of these solutions with chitosan alone can be electrospun into fibers. Water-soluble chitosan derivatives have been synthesized but could not be electrospun from aqueous solutions alone. Carboxymethyl chitosan (CMCS) in varied molecular weights (Mv = 40 to 405 kDa) and degrees of substitution (DS = 0.25 to 1.19) has been synthesized by alkalization of chitosan, followed by carboxylmethylation with monochloroacetic acid to show excellent water solubility, however, electrospinning aq. CMCS requires the addition of another fiber-forming water-soluble polymers, such as PEO, PAA, PAAm and PVA [22]. Other water-soluble chitosan derivatives, poly (ethyelene glycol) (PEG, Mn = 500 Da to 2 kDa) grafted chitosan (Mv = 137 to 400 kDa), i.e., PEG-N-chitosan and PEG-N,Ochitosan, have been synthesized via reductive amination and acylation, respectively. However, none of these aqueous solutions on their own could be electrospun into fibers [23]. PEG-N,O-chitosan is, however, soluble in organic solvents, including CHCl3, DMF, DMSO and THF, and can be electrospun by using a cosolvent to increase solution viscosity or adding a non-ionic surfactantto reduce solution surface tensions. For instance, fibers with diameters ranging from 40 to 360 nm with average diameter of 162 nm have been electrospun from 15% PEG-N,O-chitosan in 75/25 (v/v) THF/DMF cosolvents with 0.5% Triton X-100TM [23]. An elegant approach to generate chitin or chitosan nanofibers is by electrospinning a chitin derivative and then reverting back to chitin or further hydrolyzing chitin to chitosan. A simple and robust acylation reaction can render chitin derivative soluble in most common solvents that are also conducive to processing into fibers, films and coatings. The C3 and C6 hydroxyls on chitin can be acetylated with butyric anhydride to convert chitin to dibutyryl chitin (DBC) [24]. The newly formed butyl groups on C3 and C6 not only disrupt the inter-molecular hydrogen bonds, but also confer non-polar characteristics, offering DBC robust solubility (13–19 wt%) in

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acetone, DMAc, DMF, ethanol and acetic acid, all appropriate for electrospinning. As all these, except ethanol, are also solvents for CA, DBC and CA, allowing their dissolution and mixing at any ratios. The optimal solvent system for electrospinning either DBC or CA alone as well as their mixtures at all ratios is 1:1 acetone/acetic acid, resulting in fully integrated CA/DBC hybrid fibers with 30–350 nm widths with no phase separation of the two polymers. Alkaline hydrolysis of CA/DBC nanofibers in any proportion proceeds easily to first regenerate CA back to cellulose then DBC to chitin, further hydrolysis at higher alkaline concentrations and/ or temperatures deacetylates chitin to chitosan from 80 to 90%, depending on conditions. Therefore, this derivatization strategy of chitin not only improves its solubility in organic liquids, but also enables its electrospinning alone and/or with CA in any proportion into homogeneous fibers of cellulose/chitin derivatives that can be easily hydrolyzed to cellulose/chitin and cellulose/chitosan hybrid nanofibers, presenting promising holistic polysaccharide functional materials.

4.6 Surface-Active Cellulose Nanofibers In addition to homogeneous hybrid cellulose nanofibers, other materials can also be integrated into the cellulose nanofiber structure as surface layers or sheath-core structures. For example, polysaccharide nano-films could be assembled on cellulose nanofibers via electrostatic forces [25]. By alternating polymers with opposing charges, cationic chitosan and anionic dextran sulfate can be deposited to cellulose nanofiber surface in a layerby-layer (LbL) fashion, forming multiple bilayer nanofilms, each of 6.4 to 9.0 nm thickness. This LbL approach can also be applied to construct surface bound lipase hydrolase enzymes as single outer layer [26] or alternating bilayers with Cibacron blue reactive dye ligand [27]. Such surface bound approach enhances enzyme loadings on cellulose nanofibers, extends their activities as well as enables retrieval and repetitive use in biocatalysis. The highly reactive surfaces of electrospun and hydrolyzed cellulose nanofibers are also evident by easy reactions to introduce hydrophobic grafts [28], hydrophilic PVA [29] and thermally responsive poly (N-isopropylacrylamide) [30] hydrogels as well as super-hydrophilic polyacrylic acid hydrogels for enzyme encapsulation [31] and amphiphilic polyethylene glycol (PEG) tethered lipase enzymes [32] to enhance enzyme activities and repetitive catalysis. Recycled CA from cigarette filter has also been co-axially electrospun as core and fluoropolymer as shell fibrous membranes as separators for lithium-ion battery [33].

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In summary, cellulose nanofibers can be robustly fabricated from electrospinning of CAs in a diverse array of binary solvent systems to allow addition of other biopolymers, synthetic polymers or nanomaterials to form either fully integrated or phase-separated hybrid, sheath-core and porous structures. Solvent selection is critical as demonstrated by the examples discussed. In electrospinning of either single component or bicomponent polymer systems, longer polymer chains, higher polymer concentration and lower proportion of short polymer chains all contribute to necessary polymer-polymer entanglement to support continuous polymer jets and improved fibers formation. This enhanced polymer chain interaction also increases resistance to the stretching force, leading to larger fibers. These cellulose nanofibers from electrospinning and hydrolysis of cellulose derivatives are regenerated and highly reactive for introducing new surface functionalities by chemical reactions and grafting. The pathways presented offer ways to fabricate ultra-high specific surface cellulose nanofibrous substrates and enabled more recent inquiries, ours as well as many others, into functional nanomaterials based on polysaccharides such as cellulose, chitin, chitosan, etc.

4.7 Nanocelluloses ‘Nanocelluloses’ have been used to describe the nano-scale fibrillar crystalline domains separated or derived from various native cellulose sources. Nanocelluloses vary in their dimensions depending on their origins and the processes by which they are derived. The rod-like nanocelluloses have been referred as cellulose nanowhiskers (CNWs), nanocrystalline cellulose (NCC) or cellulose nanocrystals (CNCs) whereas the longer fibrils of varied lateral dimensions are termed microfibrillated cellulose (MFCs), cellulose microfibrils, nanofibrillated cellulose (NFCs) or cellulose nanofibrils (CNFs). Small organisms such as bacteria [34], algae [35] and marine animal tunicates [36] synthesize fibrillar nanocelluloses. Among these, bacterial cellulose (BC), also referred as bacterial nanocellulose or microbial cellulose, is extensively studied. BC is synthesized by a gram-negative strain of acetic-acid-producing bacteria, the acetic bacterium  Acetobacter xylinum (or Gluconacetobacter xylinus) in a fermentation process [37]. BC is extracellular gel-like product excreted into the culture medium in the form of pellicles in a web-like network of continuous ca. 10 to 100 nm wide nanofibers. In plants, fibrillar bundles of cellulose are synthesized by a hexagonal array called rosette in the presence of hemicellulose and lignin.

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By separating and/or removing the amorphous or less ordered cellulose chains, the crystalline domains can be extracted from a broad range of natural sources. Both rod-like and fibrillar nanocelluloses are mostly produced by top-down approaches from higher plants. Nanocelluloses may be in either cellulose Iα or Iβ crystalline forms, depending on their origins. Native cellulose is synthesized in two allomorphs or crystalline forms, i.e., cellulose Iα from algal and Iβ from higher plants, as confirmed by high resolution, solid state 13C NMR spectroscopy [38]. Relatively uniformly sized CNCs can be produced from a single source under a fixed condition. However, widely varied dimensions (3–70 nm widths and 35–3,000 nm lengths) have been reported from different cellulose sources and hydrolysis conditions [39–41]. By hydrolyzing and removing cellulose chains in the less ordered regions, CNCs generated are more crystalline than the original sources but usually at less than 30% yields [40, 41]. Hydrochloric acid, hydrobromic acid as well as mixed acetic and nitric acids are also capable of hydrolyzing cellulose into CNCs, but without esterifying the surfaces as in the case with sulfuric acid, while in higher yields [44, 45]. The dimensions of CNCs have shown to be mostly 5 to 20 nm wide and 100 nm or longer, dependent on the sources of cellulose as well as the hydrolysis conditions, including temperature, time and agitation [41]. Longer CNFs can be produced by shear, high-energy forces or chemical means. Shear force processing involves repetitive processing aqueous wood cellulose suspensions or pastes using mechanical homogenizers [46], cryogenic grinders or microfluidizers, high–pressure homogenizers and ultrasonic homogenizers [47] or steam explosion [48]. Further details about mechanical processes in deriving NFC can be found in a recent review [49]. The most reported chemical method for CNF derivation involves 2,2,6,6-tetramethylpyperidine-1-oxyl (TEMPO)-mediated oxidation of C6 primary hydroxyl using nitroxyl radicals in the presence of sodium hypochlorite and sodium bromide at pH around 10 and ambient temperature [50–52] and can reach a 90% yield when aided with mechanical means [53]. Due to oxidative effect, CNFs tend to be less crystalline than the sources from which they are derived. In addition to their nanoscale lateral dimensions, nanocelluloses are generally recognized as among the strongest materials, far exceeding any of high strength synthetic polymers. The uniquely high moduli were first reported on fibrillar nanocelluloses synthesized by small organisms, such as tunicate. The elastic modulus of tunicate CNWs has been reported to be ca. 143 GPa using the Raman spectroscopic technique [54] and the elastic moduli of single microfibrils from TEMPO-oxidation and acid hydrolysis

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from tunicate were measured by AFM to be ca.145 and 150 GPa, respectively [36]. Early reviews of nanocelluloses have focused mostly on those derived from bacteria and wood cellulose [55, 56]. The tremendous interest in nanocelluloses is evident from the increasing numbers of publications. As the interest in nanocelluloses has escalated over the past decade, further advancement on nanocelluloses and their applications have been extensively reviewed [57–62]. Hence, the following discussions focus on aspects of nanocelluloses from the much less studied agricultural biomass, in particular rice straw, to offer current insight on the origin- and process-linked aspects.

4.8

Nanocelluloses from Agricultural By-Products

As cellulose is the nature’s most abundant polymer, it is found in a wide range of origins. From the source perspective, major plant fibers whose features, strength and most significant constituent that are attributed to cellulose have been classified into bast fibers (flax, jute, hemp, ramie, kenaf, etc.), leaf fibers (abaca, sisal, pineapple, etc.), seed fibers (cotton, coir, kapok, etc.), core fibers (kenaf, hemp, jute, etc.), grass and reed fibers (wheat, corn, rice, etc.) and others (wood, roots, etc.) [63]. From the perspective of human activities, cellulose may be classified as primary sources for textiles, paper and wood and biofuel [63, 64], secondary nonprocessed residues from agricultural/forestry activities and by-products of food industry, such as: bark, straw, leaves, husks as well as tertiary wastes from the use, transformation and conversion of cellulosic biomass [65]. Therefore, nanocelluloses may be viewed from a much wider perspective in terms of both origins and processes, thus the focus of the following. There has been an increasing interest in deriving nanocelluloses from more diverse, non-wood and already existing sources, such as agricultural and industrial lignocellulosic residues and byproducts. Current uses of these lignocellulosics are mainly for low value-added practices, such as feed and bedding for livestock, composting, soil fertilization or combustion for energy recovery. With the advent of biofuel production from lignocellulosic biomass, generating green chemicals and materials from these sources have also become logical solutions to save not only our limited fossil fuel resources but also to reduce the environmental impact. While biomass possesses diverse chemical compositions, structures, and properties depending on their origins, one major distinction of non-woody biomass is their lower content or absence of lignin, making it less recalcitrant

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to chemical, microbial and enzymatic digestion than woody biomass. With the exception of the fibrous crops such as cotton and flax that contain significantly higher cellulose contents, the cellulose contents in most agricultural biomass vary between one third to two thirds [66], thus are close to or higher than that in woody biomass, making them viable sources for cellulose. Furthermore, cellulose in these agricultural and industrial by-products may be less integrated in their matrixes to be potentially more easily extracted. The top-down approaches to produce nanocelluloses involve either removing amorphous cellulose chains by chemical hydrolysis or separating the crystalline domains apart from each other via chemical reactions and/ or mechanical forces using cellulose isolated from various sources of higher plants. With a recent review on nanocelluloses from various agricultural crop residues and industrial wastes [65], these following analyses specifically detail the origin and process-linked effects.

4.9

Source Effects – CNCs from Grape Skin, Tomato Peel, Rice Straw, Cotton Linter

Rod-like CNCs are typically derived by hydrolysis and removal of amorphous cellulose using various acids, most commonly sulfuric acid. Cellulose in the more accessible amorphous domains is hydrolyzed by sulfuric acid via rapid protonation of glucosidic oxygen (1) or cyclic oxygen (2) by protons from the acid followed by scission of the glucosidic bonds (Figure  4.4a) [18], fragmenting into small sugars. In due process, CNC surface hydroxyls are esterified to sulfate groups that introduce negative charges to their surfaces (Figure  4.4b), thus preventing aggregation and keep CNCs well dispersed in a stable aqueous suspension. CNCs derived from different sources under the same sulfuric acid hydrolysis (64–65% H2SO4, 45 oC) have shown very different nano-scale dimensions, reflecting the structural distinctions that stem from their origins. From the same 30 min reaction, grape skin CNCs consist of mainly nanoparticles less than 5 nm in diameters among a few 103

57.3 ± 2.7

143.3 ± 30.9

–

n.a.

n.a.

1.29

0.24

–

136

153 ± 60

497 ± 161

440 ± 127

–

77.9

81.5

64.4

90.7

CrI (%)

Self-assembled ZetaSurface Fiber potential charge width Length (nm) (mV) (mmol/g) (nm)

* out of 69.8% cellulose hydrolyzed with glucose as remaining product at 92% selectivity.

6.8

H2SO4 hydrolysis 64%, 45 °C, 45 min [67,69]

Process

Yield (%)

Nanocellulose

Table 4.2 Nanocelluloses from rice straw derived by different processes and their self-assembled fibers.

–

345

320

269

234

–

10.2

20.3

19.0

23.9

Char at 500 °C Tmax (°C) (%)

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this porous carbon nanofiber bound sulfuric acid catalyst are less than half as wide or thick as the rigid rod-like CNCs (4.7 nm thick, 6.4 nm wide and 143 nm long) produced by sulfuric acid. Not only this solid acid catalyst is biobased, easily separated from both CNF and glucose products, but also can be used repetitively making the process ‘green’. These processes, i.e., sulfuric acid, coupled TEMPO and blending, blending, ACC and solid acid catalyst, not only produce nanocelluloses of different geometries, but also surface chemistries and charges. While the strong sulfuric acid hydrolysis disintegrates and breaks the less ordered cellulose chemically to release the shorter flat rod-like CNCs, the others can be optimized to defibrillate cellulose into longer CNFs and in higher yields to approach 100% with the exception of the solid acid catalyst that, in fact, generates glucose as the dominant co-product, thus also highly efficient. These geometric and surface chemical characteristics of nanocelluloses also cause them to self-assemble into distinctively different sizes and structures exhibiting amphiphilic behaviors. For instance, the extents of surface carboxylation on CNFs may be increased by increasing the levels of oxidant in the TEMPO-mediated oxidation reaction. CNFs with increasing surface carboxylate densities of 0.59, 0.92, and 1.29 mmol/g assembled into increasingly larger fibers with respective diameters of 125, 327, and 497 nm [75]. Alternatively, the surface carboxylates on CNFs derived from the same coupled TEMPO oxidization and blending may be protonated to different levels from nearly 90% charged sodium salt form to completely uncharged carboxylic acid form to exhibit different behaviors [81]. The mostly charged CNFs assembled into finest and most uniform fibers ( = 137 nm) that are more hydrophobic than hydrophilic whereas the fully protonated and uncharged CNFs assembled extensively into porous and mostly ultra-thin film-like structures. Therefore, CNFs could be tuned by varying the degree of surface carboxylation as well as protonation to desired hydrophobic-hydrophilic characteristic along the amphiphilic scale, liquid behaviors and self-assembled fiber morphologies. With the exception of CNFs from coupled TEMPO and blending, CNFs self-assemble into more crystalline fibers than the original rice straw cellulose (CrI = 72.2 %), with that of CNC being most as expected.

4.11 Ultra-Fine Cellulose Fibers from Electrospinning and Self-Assembled Nanocellulose Nanofibers and ultrafine submicrometer wide fibers can be fabricated either with electrospinning of cellulose esters, such as CAs, then

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hydrolyzed back to cellulose or via self-assembling of nanocelluloses from either the rod-like CNCs or nanofibrillar CNFs. The electrospun cellulose fibers vary in sizes, appearing similar to and in comparable sizes as the microfibrils observed on cotton fiber cell walls (Figure  4.5), but are in cellulose II crystalline structure as opposed to the highly crystalline cellulose I in cotton fibers. Both approaches are preceded by top-down processes, i.e., cellulose dissolution for electrospinning and defibrillation for nanocelluloses, that affects down to the molecular level in the former, but only separating the crystalline domains for the latter. Dissolution not only breaks all inter-molecular hydrogen bonds but destroys the crystalline structure. However, neither approach is as far top-down as that in biorefining that breaks down lignocellulosics to the much smaller molecular biofuels and chemicals. Due to their different extents of top-down processes, the crystalline structures of cellulose nanofibers from electrospinning of CA and hydrolysis back to cellulose and those from self-assembling of nanocelluloses are cellulose II and I, respectively. As cellulose I and II have distinctively different Young’s moduli of 138 GPa and 88 GPa, respectively [82], the nanofibers self-assembled from nanocellulose are expected to be far stronger and stiffer than those by electrospinning. The electrospinning approach tends to generate larger sub-micron sized fibers with smooth surfaces whereas fibers self-assembled from nanocelluloses tend to be smaller and whose sizes are highly dependent on liquid media and drying conditions. Under certain conditions, these two approaches produce similar size fibers

Figure 4.5 Electrospun cellulose nanofibers (inset) and cotton fibers (background).

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(Figure 4.6), but with drastically different chemical and physical attributes. Fibers formed by electrospinning are generally packed in very thin fibrous sheet or mat whereas those from nanocelluloses may be in any forms determined by the size and shape of containers in which they are dried. Depending on the types of nanocelluloses, the self-assembled fibers are uniquely amphiphilic, and can be tuned to be more hydrophobic by way of media in which they are dispersed and dried from. Electrospun CA fibers are hydrophobic, and become hydrophilic once the surfaces are hydrolyzed while the unhydrolyzed cores may retain some level of hydrophobicity and their thermoplastic nature and completely hydrophilic when fully hydrolyzed into cellulose II. Stemmed from their crystalline structural differences, electrospun cellulose fibers are more accessible to reagents and more chemically reactive than those assembled from nanocelluloses, lending them to be amenable for chemical modification and for chemical, biochemical and biological applications. In both approaches of electrospun and self-assembled nanofibers, the nanometer to submicron lateral dimensions contribute to their high specific surface characteristics, whereas the much higher aspect ratios offer unique oriented fibrous and porous structures with many potential applications, such as nanocomposites, reactive nanomaterials for separation membranes, medical, pharmaceutical, hygiene, absorbent and cosmetic products, energy devices and beyond.

4.12 Further Notes on Nanocellulose Applications and Nanocomposites As the emerging nanomaterial, nanocelluloses’ unique attributes have drawn and will continue to attract significant interest in many potential

10 m

(a)

(b)

Figure 4.6 Cellulose nanofibers from: (a) electrospinning of CA and alkaline hydrolysis; (b) self-assembling from freezing ( 196 °C) and lyophilization ( 50 °C) of CNFs from coupled TEMPO mediated oxidation and blending of rice straw cellulose.

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applications. For instance, rice straw nanocelluloses have been shown to exhibit unique dispersing and emulsifying properties for oil-water emulsions [83, 84], coagulants for microbes [85] as well as templates for nanoparticles and nanoprisms [86]. These nanocelluloses can also be electrospun into new nanofibrous structures [87] or assembled into super-absorbent hydrogels and amphiphilic aerogels [88] for applications such as oil recovery and separation, water purification, etc. Others have reviewed nanocellulose colloids assembled at liquid–liquid and air–liquid interfaces [89] and other related applications [90]. Related to food, nanocelluloses have been studied as stabilizing agent, food additive [91] and functional food ingredient [92] as well as additives in plastics for food packaging [93]. Nanocelluloses as part of nanocomposites is perhaps the most studied area since CNWs were first reported as reinforcement for composites [94] and has been extensively reviewed as well [95–101].

Acknowledgement The author greatly appreciates the contribution of all co-authors on their own work cited and the support for research from the National Textile Center, USDA/NIFA, NSF, NIH, California Rice Research Board and the AgTech Innovation Center at the University of California, Davis.

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5 Advanced Green Composites with High Strength and Toughness Anil N. Netravali Department of Fiber Science and Apparel Design, Cornell University, Ithaca, NY, USA

Abstract Advanced composites have been used in many applications from aerospace to sporting goods for many years because of their high specific strength and stiffness. This chapter describes the recent developments in the fabrication of advanced green composites using the newly developed liquid crystalline cellulose (LCC) fibers that exhibit high strength compared to any cellulosic fibers. The LCC fibers have high crystallinity and molecular orientation that result in strength of about 1600 MPa and Young’s modulus of around 45 GPa. Recent research efforts in combination of chemical mechanical and thermal treatments have increased both crystallinity and molecular orientation resulting in strength of around 2000 MPa and Young’s modulus of about 65 GPa. Advanced green composites made using the treated LCC fibers could have strength of up to 1 GPa and modulus of over 35 GPa. These composites could be suitable as load bearing structural components. Some of the advanced green composites also have toughness comparable to Kevlar® based composites and can be considered in ballistic applications. An exciting fact about these composites is that they are fully biodegradable and can be composted at the end of their life instead of ending in landfills, as the conventional composites do. In future, new developments in spider silk like protein fibers could also provide the required reinforcement for advanced green composites. Keywords: Green composites, advanced green composites, soy protein based resin, starch based resin, liquid crystalline cellulose fibers

Corresponding author: [email protected] Anil Netravali (ed.) Advanced Green Composites, (97–110) © 2018 Scrivener Publishing LLC

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5.1 Introduction Advanced composites based on graphite, aramid or glass fibers and epoxy resin have been used in applications from aerospace to sporting goods for many decades. Their high tensile strength and stiffness along with significantly lower density compared to metals result in high specific properties making them suitable for applications where weight reduction is critical. For example, Boeing 787, commonly known as Dreamliner, as well as Airbus 350 use airframe comprising of over 50% carbon fiber reinforced composites [1]. This has resulted in about 20% savings in the aircraft weight compared to conventional metal designs and provides significant fuel saving. Use of advanced composites also comes with other advantages for the Dreamliner and Airbus 350 such as greater cabin pressure and larger windows that provide more comfort for the passengers. The use of carbon and glass fiber reinforced composites in automobiles, though small at present, has been accelerating in recent years because of the US government regulations that require significantly higher corporate average fuel efficiency (CAFÉ) [2]. As more cars go electric and light-weighting becomes critical, the role of plastics and composites in the car can be expected to increase significantly [3]. The newest trend in advanced composites is their intended application in civil structures such as bridges and buildings. Composites can be easily molded to desired shapes instead of the conventional metal fabrication that is not only time consuming but also heavy and cumbersome. Since metals are processed at much higher temperatures than composites, they also need significantly higher amount of energy. In addition, by combining the additive manufacturing with pultrusion and, possibly, other composite forming methods, it should be possible to build larger components for structural application. Also, many parts could be manufactured at the construction site saving significant amount of transportation costs. Another inherent advantage of composites and plastics is that they do not rust like steel and hence can provide long term durability. Composites combine two dissimilar materials such as fibers and resins and require that there be excellent interaction or bonding between the two components. Good fiber/resin interfacial bonding may be achieved by fiber surface modification or by modifying the resin by using certain additives. However, good fiber/resin bonding also makes it impossible to separate the two components at the end of their life for reusing or recycling them. Composites where thermoplastic resin is used, it is possible to a limited extent, to re-form or remold the composites into different shapes and reuse them. However, in composites where the resins are thermoset such as epoxies, it is impossible to recycle or reuse the composite as the

Green Composites with High Strength and Toughness 99 cross-linking cannot be reversed. As a result, most of the composites end up in landfills at the end of their useful life although there are some efforts to recycle or reuse them. In the anaerobic conditions of the landfills the composites may not degrade for several decades or even centuries making that land useless for any other application. A small fraction of the composites are now incinerated to recover energy. In some cases, particularly involving carbon fibers, only the resin may be burnt to recover the fibers. Some studies have suggested that such recovered carbon fibers may have significantly lower strength. In any case, both landfilling and incineration are expensive, wasteful and can contribute to pollution. As the applications of the advanced composites increase in the future, particularly in mass volume applications such as automobiles and civil construction, their disposal is bound to become a serious problem. As mentioned above, no eco-friendly solution to deal with composite waste is available at present. Most fibers and resins used to make advanced composites are derived using petroleum as the raw material. While composite disposal is a significant problem, sustainability of petroleum itself is the second major problem associated with these composites. In the long run this will not be sustainable as it has been estimated that the world petroleum stock will last for another 5 to 6 decades only, at the current rate of use [4]. While bulk of the petroleum is used as fuel at present, about 6% of it is used for manufacturing various polymers/plastics, fibers, chemicals, etc. What is certain is that the current easy supplies of petroleum will end soon and we will have to search deeper and farther out in the sea for new petroleum. This would make it costlier than the renewable energy sources such as photovoltaic (solar) and wind.

5.2 ‘Greener’ Composites Many researchers have used materials derived from plants as reinforcement or filler in synthetic polymers/resins to make ‘greener’ composites [2]. Plant-based fibers and particles tend to be less expensive and abundantly available in all parts of the world. As a result, their popularity as filler or reinforcement has grown over the past few years. They also have many other advantages such as being non-abrasive and CO2 neutral when burned. In addition, the cellular and hollow nature of some fibers makes them excellent acoustic and thermal insulators. Plant-based fibers have been processed using conventional methods for centuries for applications such as apparel, furnishing, packaging, etc., and hence no new equipment is needed. They are also nonhazardous to humans and animals.

100 Advanced Green Composites Many examples can be found where thermoplastic polymers have been reinforced with plant-based fibers or particles [5]. Wood flour, saw dust and crushed wood particle based polypropylene (PP) or polyvinyl chloride (PVC) resin composites, commonly called ‘wood polymer composites’ (WPCs), can now be found in decking, window or door framing, outdoor furniture and even railroad ties. They are also referred to as ‘plastic lumber’ since they are used in the same applications where wood is commonly used. Such WPC can be easily extruded in desirable sizes. They have mechanical properties comparable to wood but they can last much longer than wood for the intended applications because of the hydrophobic characteristics of the polymers used. Another advantage of WPC is that they can be made using recycled plastics rather than virgin plastics as well as used wood products. While WPC is commonly pultruded in the form of planks, they can also be molded into different shapes which makes their processing much easier compared to wood which needs complex processes to shape them. Since the starting wooden log could be large, getting it to desired shape can generate significant amount of waste. It is well known that higher aspect ratio of the reinforcing component, for example, fibers, can result in higher tensile properties of the composites. Since fibers have much higher aspect ratio compared to saw dust or wood particles, stronger, stiffer and tougher composites can be obtained by using natural fibers as reinforcing elements. However, no plant-based fibers are continuous and are obtained as short length or ‘staple’ fibers only. They can only be used as twisted yarns or converted into woven or knitted fabrics or nonwoven webs. Fibers such as flax, henequen, hemp, ramie, sisal, abaca, etc., have good tensile properties and those properties get reflected in the composite mechanical properties. Ordinarily the strength of these fibers varies between 200 and 600 MPa and modulus varies between 10 and about 50 GPa, though some can be stronger and stiffer. Tensile and physical properties of many natural fibers can be found elsewhere [2, 6]. When they are spun into yarns, to obtain continuous form, their strength and modulus drop significantly as a result of the twist. This automatically results in lower composite properties. Use of woven and knitted fabrics and nonwoven webs also result in composite strength reduction. However, these fiber forms allow fabricating laminates that may be organized to form layered composites with desired properties in different directions. Many examples can be found in literature where plant-based fibers in woven, knitted or nonwoven forms have been combined with conventional thermoplastic or thermoset resins that are non-degradable [7–12]. Since these greener composites combine petroleum based non-degradable resins with degradable fibers, they can neither return to an industrial metabolism

Green Composites with High Strength and Toughness 101 nor to a natural metabolism [2]. That means they cannot be food stock for either of the systems. They can only be downcycled as their properties degrade when reprocessed or put in landfills [2]. They may also be incinerated to recover energy. However, as mentioned earlier, incineration can be expensive and can result in significant pollution.

5.3 Fully ‘Green’ Composites As a result of severe environmental concerns posed by the non-degradable composites and fast depletion of petroleum reserves the current trend is to fabricate a new class of fully green composites that use both fibers and resins that are biodegradable. These green composites can be easily disposed of or composted without using any landfill capacity. Most green composites are also made using plant-based fibers and resins that are fully sustainable, making them ‘green’ in every sense [13–18]. While this research is in its infancy at present, the number of research papers have been increasing almost exponentially for the past 10 or so years and a few products have already entered the market [2]. They can be an excellent choice for consumer products such as ‘disposables’, intended for one-time use, or those that have short life cycles or ‘nondurables’. However, when they are used indoors avoiding contact with water, they can last for several years or decades without degrading. It is intuitive that composite properties are a function of the properties of its constituents, that is, the fibers and the resins that are used to fabricate them. However, the strength and stiffness of the composites are determined primarily by the properties of the fibers, the reinforcing elements used in composites, as well as their volume content and orientation. Since most green composites are made using plant-based fibers they have limitations in terms of their mechanical properties. Even though some references may be found out where the plant-based fibers have been shown to have strengths close to 1000 MPa or even higher, most commercial plantbased fibers have strength below 600 MPa and modulus below 50 GPa [2]. As mentioned earlier, all plant-based fibers are short and have to be twisted into yarns to obtain a continuous form and be woven or knitted into fabrics. Twisting fibers can reduce their effective strength and modulus even further. For composites that contain about 50% fibers by weight, this can translate into moderate strength between 250 and 300 MPa and modulus between 20 and 25 GPa. These values are somewhat comparable to wood and hence these composites can be used where wood or wood based materials such as particle board, medium density fiberboard (MDF) or plywood are currently used.

102 Advanced Green Composites

5.4 ‘Advanced Green Composites’ As mentioned above, the tensile properties of composites are primarily a function of the fiber properties. However, fiber/resin interface also plays a significant role. As the fibers in the composites begin to break under tensile stress, the resin transfers the load from broken fibers to intact fibers through the fiber/resin interface [19]. This results in enhanced strength and stiffness of the composite. Thus, it is critical to have good fiber/resin bond if high strength and stiffness are desired. However, strong fiber/resin interfacial bonding also makes composites brittle. A weaker fiber/resin interface, on the other hand, promotes interfacial debonding mechanism that allows energy absorption, making the composites tougher. In summary, for stronger composites it is critical to use stronger fibers and to have good fiber/ resin interfacial bonding as well. It is known that none of the plant-based natural fibers are as strong as some of the synthetic fibers such as carbon, Kevlar®, ultra-high molecular weight polyethylene (UHMWPE), or glass fibers used in conventional advanced composites. Natural fibers lack the molecular orientation and crystallinity of the synthetic fibers such as Kevlar® or UHMWPE which lowers their strength and stiffness. In addition, there can be many defects that reduce their strength even further. While the defects in plant-based fibers may not be rectified, chemical and/or mechanical treatments can be used to enhance their mechanical properties. In general, molecular orientation and crystallinity of these fibers can be improved by treating the fibers under tension. Such morphological changes directly result in improvement in their strength and/or stiffness. However, such treatments carried out under slack conditions can improve the toughness of the fibers by increasing their fracture strain. These treatments on natural cellulosic fibers have given a great opportunity to improve their tensile performance [20, 21]. A variety of chemical modifications including acetylation and alkali treatments and their effects on natural fibers have been studied in the past [11, 22, 23]. Gomes et al. did a systematic study of properties of curaua fibers and starch based green composites [23, 24]. They used highly concentrated alkali (NaOH) to treat curaua fibers. The alkali treated curaua fibers not only showed significant improvements in their tensile strength but the fiber/starch interfacial bond strength also improved as a result of removing the hemicellulose and lignin [25, 26]. They also explored 3 different methods of composite fabrication; direct placing of fibers (in the mold), pre-forming (of tapes) method and Prepreg sheet method and found that the fabrication methods that affect the fiber orientation, cannot give good composite properties. Their results showed strength of the composites

Green Composites with High Strength and Toughness 103 increased from 216 MPa for direct method to 327 MPa, over 50% increase, for prepreg method. Even more impressive was the increase in the composite modulus which increased from 13 (for untreated fibers) to 36 GPa for the respective methods. As expected, the fracture strain decreased, from 1.53% to 1.16%. The alkali treated fibers resulted in tougher composites as well. Compared to many chemical treatments available, alkali treatment may be the least expensive and environment-friendly way to improve the mechanical and interfacial properties of cellulosic fibers. It does not use any toxic organic chemicals and the alkali can be easily neutralized after the treatment. Goda et al. also studied the effect of load during mercerization on ramie fibers and their composites [25]. Mercerization, a treatment involving NaOH, is commonly used for cotton yarns and fabrics for improving their tensile and other properties. The treatment has been shown to increase the cellulose content through removing the hemicellulose and the lignin [26]. Goda et al. also found that application of load during mercerization treatment increased the fiber strength significantly compared to control (untreated) fibers and that the increase was a function of the load applied [25]. In their study the highest strength obtained for treated fibers was 661 MPa. It was hypothesized that the improvements in the fiber properties were a result of the morphological changes such as molecular orientation, microfibrillar angle as well as improved crystalline alignment along the fiber axis [27]. The increased properties were reflected in composite properties as well. Application of mercerization and related morphological changes and corresponding enhancement in mechanical properties has been observed by other researchers as well [27–29]. Kim and Netravali mercerized sisal fibers, both under tension and without tension, and fabricated unidirectional green composites using soy protein based resin [30]. They also found that both strength and Young’s modulus of sisal fibers increased for both mercerization treatments, with and without the tension. However, fibers treated under tension showed much higher fracture stress of 382 MPa compared to control (284 MPa) and slack treated fibers (339 MPa). Young’s modulus of the fibers treated under tension was significantly higher at 11 GPA compared to control (5.24 GPa) and slack treated fibers (6.12 GPa). The tensile property changes were in agreement with the suggested improvement in the fiber morphology by other researchers. The strength and modulus of fibers were reflected in the unidirectional green composites made using soy protein as the resin [30]. Patil et al. used a combination of mercerization followed by a heat treatment regime to improve sisal fiber tensile properties [31]. Both mercerization and heat treatment on sisal fibers were carried out under tension. The treatment significantly improved the Young’s modulus of the fibers

104 Advanced Green Composites from 5.5 to 16.7 GPa, over 200% increase, and strength from 300 to over 450 MPa, an increase of 50%. They used non-edible protein derived from karanja (Pongamia pinnata) and starch from mango (Mangifera indica) seed kernel, agricultural waste products, to fabricate unidirectional green composites. The strength of both mango seed kernel starch and karanja seed protein based composites increased by about 50% and their Young’s modulus increased by about 80% in the case of composites using alkali/ heat treated fibers compared to control fibers. In all the examples presented in the paragraphs above, green composites with higher strength were made using modified natural fibers. However, even with the modifications through chemical and mechanical treatments the natural fiber strength and Young’s modulus are no match for the conventional high strength fibers such as Kevlar® or glass and the resulting green composites reinforced with these modified natural fibers cannot have high strength or Young’s modulus to be called ‘advanced’ in its true sense. To obtain high strength and stiffness it is critical to have fibers that have desired high strength. Recent development of high strength LCC fibers obtained by dissolving cellulose in phosphoric acid and spun using air gap wet spinning technique, similar to that used to spin Kevlar® fibers, has created a significant opportunity to develop advanced green composites with high tensile properties [32–34]. Netravali, et al., used the LCC fibers to fabricate high strength composites with soy protein based resins which were termed as ‘advanced green composites’ because of their high tensile properties [35]. They used interpenetrating network-like (IPN-like) soy protein based resin, with improved tensile properties. The IPN-like soy protein based resin was formed by the addition of a polycarboxylic acid which cross-links by itself. They further added microfibrillated cellulose to the resin to improve the properties even further. The strength and Young’s modulus of the LCC fibers they used were found to be 1680 MPa and 40 GPa, respectively. This strength was 3 to 4 times higher than any commonly used natural fibers. Another significant advantage of the LCC fibers was that they came in continuous form and did not need any twist as in the case of natural fiber yarns. Twist is known to lower the strength and Young’s modulus of the yarn compared to the fibers used because of the obliquity effect. The continuous form of LCC fibers also makes it convenient to use conventional industrial equipment to fabricate composites. The strength of the unidirectional LCC fiber reinforced composite was 638 MPa, highest ever obtained for any green composites, with only 40% fiber volume fraction. At 65% fiber volume fraction, normal for any commercial composites, these green composites would have strength of over 1 GPa, making them truly ‘advanced green composites’. These composites

Green Composites with High Strength and Toughness 105 were made using hand lay-up which does not result in uniform tension or good fiber orientation. If commercial machines are used, much better fiber orientation and tension can be maintained and, hence, much better composite properties can be obtained. The authors found the flexural properties of these advanced green composites to be high as well. They also fabricated composites using Kevlar® and glass fibers, separately, using the same soy protein based resin and compared the toughness of the advanced green composites with them and found that the fracture toughness of the advanced green composites to be much higher than both Kevlar® and glass fiber reinforced composites. Based on the toughness results of their static tests they concluded that these composites may also be used in ballistic applications. Kim and Netravali treated similar LCC fibers with various alkalis and found that KOH, milder than NaOH, worked well to improve their tensile properties [36]. They treated the LCC fibers with 2M KOH solution at room temperature (RT) under no load (slack) as well as under a predetermined (7% of the fracture stress) load for 2 h. The resin used in their study was prepared using soy protein isolate (SPI). They observed no change in the fiber surface after the KOH treatment. The X-ray diffraction (XRD) results indicated a significant increase in the fiber crystallinity after the KOH treatment; slack or under load. As a result, both tensile strength and Young’s modulus showed significant increases. For example, strength increased from 1483 to 1588 MPa (7% increase) for fibers treated in slack condition and 1749 MPa (about 18% increase) for the fibers treated under load. Similarly the Young’s modulus increased from 47.8 to 53.7 GPa for the slack treated fibers and 63.7 GPa (over 33% increase) for the fibers treated under tension but fracture strain decreased from 6.8% to 6.5% and 5.2%, respectively. The LCC fiber/SPI green composites had a strength of 540 MPa (untreated LCC fiber), 583 MPa (slack treated LCC fibers) and 652 MPa (LCC fibers treated under load) and Young’s moduli of 18.7 GPa, 20.1 and 24.1 GPa, respectively. These composites had a fiber volume fraction of only about 41.5%. They also concluded that with 65% fiber volume fraction the Young’s modulus of the composites fabricated with LCC fibers treated under load, would be 37 GPa and the strength would be over 1 GPa, making them advanced green composites in every sense. Recently, LCC fibers were modified using a combination of chemical and heat treatment, under a predetermined tension [37]. The chemical treatment was carried out using 5% NaHSO3 solution at RT under a tension equivalent to 7% of the single LCC fiber strength for a period of 1  h. After washing with DI water, the fibers were dried at 140 ºC in an air circulatory oven for 1 h under same tension. The surface topography,

106 Advanced Green Composites crystallinity, molecular orientation and tensile properties were investigated to assess the changes after the treatment. The strength, Young’s modulus and fracture strain of the treated LCC fibers were 1995 MPa (33% increase from 1498 MPa for control fibers), 68 GPa (41% increase compared to 48 GPa for control fibers) and 5.2% (15% decrease compared to 6.2% for control fibers), respectively. Such high properties of cellulose fibers have not been reported in open literature. Rahman and Netravali also fabricated advanced composites by combining waxy maize starch (WMS) based resin with control and treated LCC fibers [37]. The WMS resin in this study was cross-linked using 1,2,3,4-butane tetracarboxylic acid (BTCA) to improve its tensile properties. In addition, micro-fibrillated cellulose (MFC) was dispersed in the WMS resin to further improve its mechanical properties. The advanced green composites fabricated using treated LCC fibers and WMS resin showed significant increase, compared to control LCC fibers. The Young’s modulus increased to 32 GPa (compared to 22 GPa for composites using control fibers) and tensile strength increased to 790 MPa (compared to 505 MPa for composites using control fibers). All composites had a fiber volume fraction of about 55%. These advanced green composites have the highest strength obtained thus far and when their density of less than 1.5  g/cc is taken into account, their specific strength would be higher than some of the commercial glass fiber-reinforced composites. When compared to steel, the advanced green composites are four to six times stronger than most varieties of steel, based on their specific (normalized per unit weight) strength. In addition to their excellent mechanical properties, the fire performance of the soy protein based resins and composites have been shown to be comparable or better than some commercially available synthetic resins [38, 39]. Other proteins such as whey protein, casein and keratin also have good fire resistance [40–42]. Starch also has been shown to have good fire resistance [43]. Composites based on these resins can be used where superior fire performance is needed.

5.5 Conclusions Advanced green composites with high strength and stiffness is truly a new paradigm within the emerging field of green composites. With such excellent mechanical properties advanced green composites could provide many new opportunities for their use in primary structural applications in the future. While at present, high mechanical properties have been possible because of the newly developed LCC fibers, opportunities could also come

Green Composites with High Strength and Toughness 107 along for using high strength protein based fibers in the future. For example, dragline silk obtained from the golden orb spider has strength in the range of 1 GPa. However, it is impossible to produce commercial quantities of the dragline silk using spiders for a variety of reasons [2]. As a result, significant research has been and still being conducted to obtain protein with spider silk-like chemistry and spin it into strong fibers. Earlier Nexia Biotechnologies, Inc., a Canada based company, was able to transfer spider gene into goats [44, 45]. This technique allowed spider silk-like protein in goat’s milk, ready for extraction. However, the strength of these fibers was not high enough. Recently, researchers at Bolt Threads, a California based company, have successfully engineered the right sequence protein using yeast, rather than using spider gene [46]. Their process can produce large quantities of this protein, needed for commercial scale, through fermentation. The obtained protein is then wet spun into their trademarked Engineered Silk® fibers. It needs to be seen if Bolt Threads or anyone else can make protein fibers that can compare with the tensile properties of LCC fibers, good for high strength composites. In addition, bacterial cellulose nanofibers which have high crystallinity and molecular orientation and show high tensile properties would also be a good candidate if easy ways to orient them can be found. Since both constituents of green composites or advanced green composites are fully biodegradable, at the end of their useful life they can be easily discarded to biodegrade in normal environment such as a compost without harming the environment or the earth. This is quite unlike the petroleum based conventional composites that mostly end up in landfills and stay there in ‘as is’ condition for several decades or even centuries, making that land useless for any other use. The most important shortcoming of the current green composites is perhaps their sensitivity to moisture. Since both cellulosic fibers and protein or starch based resins are hydrophilic, green composites made using them, do absorb water and swell when in contact with high humidity environment or water. As they absorb water and get plasticized, their mechanical properties are affected. While this aspect could discourage their use, protecting them from water should not be difficult. Several hydrophobic and ultra-hydrophobic paints, coatings and varnishes, some produced using soybean oil, are commercially available and can be used to protect green composites including advanced green composites. It should also be possible to prepare hybrid composites where a very thin protective shell is made up of conventional composite (e.g., glass or graphite and epoxy) covering the bulk of the advanced green composite core that provides majority of the structural and load carrying capability. Such conventional

108 Advanced Green Composites composite shell can not only protect the advanced green composites from water or other solvents, but also provide additional load carrying capability.

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6 All-Cellulose (Cellulose–Cellulose) Green Composites Shuji Fujisawa1,2*, Tsuguyuki Saito2, and Akira Isogai2 1

Department of Forest Resource Chemistry, Forestry and Forest Products Research Institute, Tsukuba, Ibaraki, Japan 2 Department of Biomaterials Sciences, Graduate School of Agricultural and Life Sciences, The University of Tokyo, Tokyo, Japan

Abstract All-cellulose composites (ACC), which are entirely bio-based composite materials, satisfy the increasing demand for environmentally friendly and sustainable materials. Because of their excellent mechanical properties, ACCs are considered a green alternative to glass- and carbon-fiber-reinforced polymer composites. This chapter summarizes progress in ACC preparation and discusses their material properties such as optical, mechanical, and gas barrier properties. Finally, we give a future perspective on ACC development for applications in various fields, including optical devices, food, and medical. Keywords: All-cellulose composite, biocomposite, green composite, selfreinforcing material, mechanical properties

6.1 Introduction 6.1.1

Cellulose

Cellulose, (1 4)-β-d-glucan, is the most abundant carbon resource in nature (Figure 6.1). It is biosynthesized by plants, bacteria, and some marine animals, and its total annual biomass production amounts to 1.5 × 1012 tons [1]. Because of its abundance and its favorable properties such as hydrophilicity, chemical stability, biodegradability, biocompatibility, and renewability,

*Corresponding author: [email protected] Anil Netravali (ed.) Advanced Green Composites, (111–134) © 2018 Scrivener Publishing LLC

111

112 Advanced Green Composites HO HO OH

OH

O O HO

HO

O O

OH HO Glucose unit

H

OH HO

n−2

O

HO

Figure 6.1 Chemical structure of cellulose (n: degree of polymerization).

cellulose has been used commercially in various materials such as textiles, films, pulp, and paper. In recent years, there has been an increasing demand for sustainable and environmentally friendly materials, and cellulose is a prime candidate for replacing many of the oil-based materials. Cellulose has a unique crystalline morphology; cellulose chains are spontaneously assembled into crystalline nanofibrils, which are referred to as cellulose microfibrils, via biosynthesis [2]. Because of their high crystallinity, the microfibrils have a high elastic modulus (130‒150 GPa) [3–5] and strength (2‒6 GPa) [6, 7], and low coefficient of thermal expansion (4‒6 ppm K‒1) [8, 9] along the c-axis direction. The aspect ratios (length/ width) of the microfibrils can be as high as over 300, with their widths ranging between 2 and 20 nm. These sizes depend on the source from which cellulose is obtained. Figure 6.2 gives an example of cellulose microfibrils, showing the hierarchical structure of wood cellulose [10]. The microfibrils present in the cell walls of wood have an ultrafine width of ~3 nm, consisting of approximately 30–40 cellulose chains, and the microfibrils play a significant role in the structural support of plant bodies. The inherent reinforcing potential of cellulose microfibrils has generated considerable interest in their use as nanofillers in polymer nanocomposite materials.

6.1.2 Nanocelluloses for Polymer Composite Materials Nanocelluloses are nanosized fibrous materials composed of isolated or consisting of a few bundles of cellulose nanofibrils. They are prepared by mechanical disintegration of native celluloses with or without chemical treatment. The size and shape of the nanocelluloses can be controlled by changing the preparation process and/or the cellulose source. Nanocelluloses prepared simply by harsh mechanical treatments such as high pressure homogenization [11, 12] or grinding [13, 14] are commonly termed as microfibrillated celluloses (MFCs). The widths and lengths of MFCs are typically in 20‒40 nm range and several micrometers, respectively.

All-Cellulose (Cellulose–Cellulose) Green Composites Tree

Fibres

Cellulose fibre

Cellulose Bundle of cellulose fibrils microfibril

113

Cellulose molecule

Fiber surface

Wood

Wood tissue

Width 20–30 m Length 1–3 mm

Width > 15–30 nm

Width 0.4 nm Width ~3 nm Length >2 m Length ~500 nm Crystallinity 70–90%

Figure 6.2 Hierarchical structure of wood cellulose and characteristics of cellulose microfibrils. This figure is based on a figure in Ref. [10] and is produced with permission from Royal Society of Chemistry (© RSC, 2013).

Such MFCs form strong and web-like network structures. Cellulose nanocrystals (CNCs) or nanocrystalline celluloses, which are also known as cellulose nanowhiskers [15, 16], can be prepared by sulfuric acid treatment followed by ultrasonication. CNCs prepared from wood celluloses have widths of ~5 nm and lengths ranging from 50 to 150 nm, and show good dispersibilities in water and some polar organic solvents. Catalytic pretreatment using 2,2,6,6-tetramethylpiperidiniyl-1-oxyl (TEMPO) introduces carboxyl groups on the microfibril surfaces. This helps nanocellulose pretreated by oxidation (TEMPO-oxidized cellulose nanofibrils, TOCNs) to be isolated at the individual microfibril level because of the repulsive forces arising from the surface carboxyl groups [10]. Nanocelluloses have attracted attention as nanofiller materials in polymer based composites because of their excellent mechanical properties and reinforcing potential resulting from their highly crystalline nanostructure. Since the first reports by Favier et al. in the 1990s [17, 18], there have been several studies of nanocellulose/polymer composite materials [19–21]. These nanocomposites typically have improved the mechanical properties and thermal dimensional stability as a result of reinforcement with nanocelluloses which offer the possibility of renewable nanofillers for sustainable industrial use.

114 Advanced Green Composites

6.1.3 All-Cellulose Composites All-cellulose composites (ACCs) have attracted significant interest in recent decades [22]. In an ACC, the reinforcing and matrix phases are both cellulose. The relatively soft regenerated cellulose phase (matrix) is reinforced by crystalline cellulose fibrils (filler). The material is therefore not only fully renewable and potentially biocompatible, but also has good mechanical properties because of the ideal fibril/matrix interfacial compatibility. This self-reinforcement system is a common concept for some semicrystalline synthetic polymers such as polyethylene [23, 24] and polypropylene [25, 26]. The thermal and mechanical properties of self-reinforced materials are better than those of amorphous ones. Nishino et al. first reported the ACC concept in 2003/2004, and since then a significant number of papers on such materials have been published [27, 28]. This chapter reviews recent advances in ACCs. First, the methods for ACC preparation are summarized and the latter part discusses their properties, mainly focusing on the mechanical properties, and evaluates the potential of ACCs as green nanocomposite materials.

6.2 Preparation of ACCs 6.2.1 Dissolution of Cellulose Cellulose dissolution is one among the most important processes in the development of ACCs because cellulose does not melt below its thermal degradation temperature. Cellulose is insoluble in water and most organic solvents because numerous intra- and inter-molecular hydrogen bonds prevent solvent molecules from penetrating between the cellulose chains. Significant efforts have, therefore, been made to find solvents for dissolving and processing cellulose materials.

6.2.1.1 Aqueous Solvents Since the discovery that cellulose dissolves in cuprammonium hydroxide aqueous solution [29], several aqueous cellulose solvents have been reported and some of which have been commercially available for decades. A representative aqueous system is NaOH/CS2/H2O, in which cellulose dissolves by forming unstable derivatives with CS2 [30]. Extrusion of the solution in acidic media regenerates cellulose by releasing the substituents. This method has been commonly used to obtain viscose rayon fibers.

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N-Methylmorpholine N-oxide (NMMO)/H2O is another solvent that is used industrially to prepare lyocell fibers, also known as Tencel® [31, 32]. Sulfuric acid or phosphoric acid based aqueous solutions dissolve cellulose at specific concentrations, accompanied by depolymerization of the cellulose molecular chains by acid hydrolysis of glycoside bonds. An environmentally friendly and low-cost aqueous NaOH system has been presented which completely dissolves cellulose at low temperatures [33, 34]. This is a simple but efficient dissolution process and the regenerated celluloses have either cellulose II or amorphous crystalline structures [34]. Addition of urea to the aqueous NaOH system significantly enhances dissolution [35, 36], and, based on this system, LiOH/urea has also been developed successfully as a cellulose solvent [37].

6.2.1.2

Organic Solvents

Dissolution of cellulose using organic solvents has both industrial and academic interest for structural analysis, chemical modification, and hybrid material fabrication. LiCl/N,N-dimethylacetamide (DMAc) has been shown to be a good solvent for cellulose [38, 39]. Also, it has been speculated that LiCl and DMAc form complex structures with cellulose molecules based on hydrogen bonds, and cellulose from various sources can be stably dissolved without significant depolymerization [40]. This system is suitable for basic analyses [41, 42] or chemical modification [43] of cellulose. Other polar aprotic organic solvent/LiCl systems such as LiCl/1-methyl-2-pyrrolidinone [39], LiCl/dimethyl sulfoxide [44], and LiCl/1,3-dimethyl-2-imidazolidinone [45] have also been used as good cellulose solvents. It has been reported that an SO2/amine/dimethyl sulfoxide system dissolves cellulose well by forming SO2/amine/cellulose complexes in the solvent [46].

6.2.1.3 Ionic Liquids Ionic liquids (ILs) have also attracted significant interest as green alternatives to conventional solvents because of their reusability, high chemical stability, and low volatility. ILs are salts that show unique solvent capabilities and are commonly defined as salts with melting temperatures below 100 °C. Graenacher found that in the presence of nitrogen-containing bases molten N-ethylpyridinium chloride acts as an efficient cellulose solvent [47]. In 2002, Swatloski et al. reported that ILs containing 1-butyl-3-methylimidazolium cations effectively dissolved cellulose in high concentrations after heating or microwave treatment [48]. Since then,

116 Advanced Green Composites cellulose dissolution using ILs has been extensively studied in applications such as chemical modification [49–51] and processing fibers, films, and composite materials [52–56]. Many reports have confirmed that noncovalent interactions between cellulose and the anions of ILs play a significant role in the dissolution process [57]. Since numerous potential combinations of cations and anions are possible, ILs can provide new opportunities for cellulose processing [57–59].

6.2.2 Preparation of ACCs Advances in cellulose dissolution have enabled significant progress to be made in ACC processing. A variety of methods have been developed and can be classified into two types; one-phase preparation and two-phase preparation as described below.

6.2.2.1 One-Phase Preparation This process includes partial dissolution of native celluloses using good cellulose solvents followed by regeneration by exposing the cellulose to poor solvents such as water, methanol, and ethanol (Figure 6.3) [60, 61]. The major advantage of this method is that interface-free composites can be prepared. Nano-welding of the two fibers that occurs in this process can effectively merge the adjacent surfaces of the original celluloses, resulting in ideal interactions via regenerated cellulose bonds [62]. The undissolved fiber content (i.e., filler content) of the composite can be controlled simply by changing the dissolution time.

6.2.2.2 Two-Phase Preparation This preparation process involves mixing two cellulose phases, i.e., solid (filler) and solution (matrix) phases. Nishino et al. were the first to propose this method [27, 28]. They impregnated dried ramie fibers in a kraft pulp solution prepared using 8 wt% LiCl/DMAc followed by regeneration and washing with methanol to remove the solvent [27, 28]. The material had a composite structure consisting of reinforcing ramie fibers (filler) surrounded by regenerated kraft pulp (matrix). Similarly, homogeneous mixtures of nanocellulose/regenerated cellulose can be prepared by mixing a cellulose solution and a nanocellulose dispersion (Figure 6.4) [63]. In such cases ACCs are formed by subsequent regeneration or solvent casting [64, 65]. One advantage of this method is that the fiber content of the

All-Cellulose (Cellulose–Cellulose) Green Composites

Cellulose fibres

(a)

Surface selective dissolution

117

All-cellulose composites

1h

2h

3h

4h

5h

6h

9h

12 h

10 m (b)

Figure 6.3 (a) One-phase preparation of all-cellulose composites (ACCs). (b) Scanning electron micrographs of cross-sections of ACCs with different immersion times. Reproduction of this figure, modified from Ref. [60], with permission from Elsevier (© Elsevier, 2008).

Activated MCC

LiCI/DMAc solution

Cellulose solution

Stirring

Drying

Composite film CNFs/cellulose Sonication + homogenization

Softwood pulp

CNFs dispersion

Figure 6.4 Two-phase preparation of all-cellulose composites. Reproduction of this figure, modified from Ref. [63], with permission from Elsevier (© Elsevier, 2014).

ACC can be easily controlled by changing the mixing ratio. Good dispersion of the reinforcement phase is achieved by controlling the dispersibility of the original nanocellulose. Table 6.1 summarizes the ACC preparation methods.

LiCl/DMAc LiCl/DMAc LiCl/DMAc LiCl/DMAc LiCl/DMAc IL (BmimCl) IL (BmimCl) IL (AmimCl) Solvent

Ramie fiber

Canola straw bleached pulp

Filter paper

Bacterial cellulose

Sugarcane bagasse pulp (MFC)

Canola straw bleached pulp

MFC

MCC

Filler Softwood pulp Cotton linter Cotton linter Cotton linter

Wood pulp (TOCN)

Ramie fiber

Tunicate, Cotton (CNC)

Cotton (CNC)

NaOH/urea/H2O

NaOH/urea/H2O

NaOH/urea/H2O

NaOH/urea/H2O

LiCl/DMAc

Avicel® (MCC)

Two-phase

LiCl/DMAc

Beech pulp

Matrix

LiCl/DMAc

Cotton linter (MCC)

One-phase

Solvent

Cellulose source

Preparation

Table 6.1 Summary of all-cellulose composite preparation methods.

H2SO4/water

H2SO4/water

H2SO4/water

H2SO4/water

Regeneration

Water

Water

Methanol

Ethanol

Methanol

Methanol

Methanol

Methanol

Water

Water

Water

Regeneration

Qi et al. [77]

Pullawan et al. [76]

Yang et al. [75]

Yang et al. [74]

Reference

Zhang et al. [73]

Shakeri et al. [72]

Yousefi et al. [62]

Ghaderi et al. [71]

Soykeabkaew et al. [70]

Nishino and Arimoto [69]

Yousefi et al. [68]

Soykeabkaew et al. [60]

Duchemin et al. [67]

Gindl et al.[66]

Gindl and Kecks [61]

Reference

118 Advanced Green Composites

Softwood pulp Avicel® (MCC) Cotton Linter MCC Avicel® (MCC) Avicel® (MCC) Avicel® (MCC) Ramie fiber Filter paper

Ramie fiber

Avicel® (MCC)

Wood pulp (TOCN)

Wood pulp (MFC)

Tunicate, Cotton (CNC)

Tunicate

Avicel® (MCC)

Ramie fiber

Rice husk

IL (BmimCl)

LiCl/DMAc

LiCl/DMAc

LiCl/DMAc

LiCl/DMAc

LiCl/DMAc

LiCl/DMAc

LiCl/DMAc

LiCl/DMAc

NMMO/H2O

Water

Methanol

Water

Water

Water

Water

Water

Water

Methanol

Water/ethanol

Zhao et al. [82]

Qin et al. [81]

Pullawan et al. [80]

Pullawan [79]

Pullawan et al. [76]

Zhao et al. [63]

Fujisawa et al. [65]

Gindl and Keckes [64]

Nishino et al. [27]

Quajai and Shanks [78]

MCC: microcrystalline cellulose. MFC: microfibrillated cellulose. CNC: cellulose nanocrystal. TOCN: TEMPO-oxidized cellulose nanofibril.

AmimCl: 1-ally-3-methylimidazolium chloride. BmimCl: 1-butyl-3-methylimidazolium chloride.

Hemp fiber

Hemp fiber

All-Cellulose (Cellulose–Cellulose) Green Composites 119

120 Advanced Green Composites

6.3 Structures and Properties of ACCs 6.3.1 Optical Properties ACCs generally have high optical transparency, irrespective of the cellulose starting material or the preparation process used. Yousefi et al. used an IL to prepare transparent ACC films with light transmittances of 76% at 800 nm (Figure 6.5) [62]. Yang et al. fabricated flexible and transparent ACC films from TOCNs and regenerated cellulose using the NaOH/urea/ H2O system; the films had high transparencies, i.e., 88% (600 nm) [74]. This high transparency was a result of the thin width of the fillers and the small difference between the refractive indices of the fillers and the matrix. ACC films, therefore, typically show high transmittances, >75 %, in the visible region [61, 63, 73].

6.3.2 Mechanical Properties In general, the elastic modulus and strength of an ACC increase proportionately with increasing filler fraction because of reinforcement provided by the filler phase. The mechanical properties of ACCs are generally better than those of most nanocellulose/synthetic polymer nanocomposites, probably because of the ideal interfacial interactions between the filler and matrix phases which are chemically identical [60]. Gindl et al. compared the mechanical properties of an ACC with those of cellulose/epoxy composites produced using the same cellulose fibers [66]. They found ACC to have a better elastic modulus and tensile strength because of good filler/ matrix adhesion [66]. Microfiber sheet

All-cellulose nanocomposite

Nanowelding

Figure 6.5 Optically transparent all-cellulose composite film. Reproduction of this figure, modified from Ref. [62], with permission from American Chemical Society (© ACS, 2011).

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Pullawan et al. investigated the stress-transfer mechanisms during ACC deformation based on well-defined Raman spectra [76]. Observation of the peak shifts of the Raman band initially located at 895 cm−1, which corresponds to the peak of cellulose II (matrix), showed that the higher the aspect ratio of the filler, the better the stress-transfer and mechanical properties of the ACC [76]. Based on the same study, it was also suggested that the mechanical properties of an ACC depend on the cellulose solvent used to prepare the matrix. The mechanical properties of ACCs prepared using a LiCl/DMAc system were better than those of ACCs prepared using a NaOH/urea/H2O system. Duchemin et al. used wide-angle X-ray scattering and solid-state NMR spectroscopy to investigate the molecular packing of cellulose chains in ACCs prepared from LiCl/DMAc [83]. In their study the cellulose matrix was found to show paracrystal properties rather than amorphous ones, suggesting that a partly ordered matrix resulted in an ACC with good mechanical properties. A finite element model has been successful in explaining the deformation and failure behaviors of both randomly and uniaxially reinforced ACCs [84, 85]. The model suggests that the elastic moduli of ACCs along the aligned axis increase with increasing aspect ratio, volume fraction, and orientation of the filler, and that the progressive failure of ACCs is a consequence of the consecutive cracks that appear during tensile testing. The force and the impact energy during deformation were evaluated in the studies, and the predictions were in good agreement with the experimental data. Because of the good filler/matrix binding in ACCs, the work of fracture, this reflects the toughness of a material, increases with increasing filler content. Yang et al. confirmed that the work of fracture increased slightly from 11 to 13 MJ m–3 with increasing filler content [74]. An ACC prepared by partial dissolution of bacterial cellulose had a fracture value as high as16 MJ m–3 [70]. However, increases in the work of fracture are not significant, and in some cases the value decreases although the modulus and strength both increase; this results in a continuous decrease in the strain at break with increasing filler fraction. Fracture depends significantly on defects in the materials, which can act as crack initiation sites during tensile testing. Duchemin et al. investigated the relationship between the structures and mechanical properties of ACCs by changing the preparation process [67]. Their results suggest that a slower preparation process can lead to a higher strain at break. The formation of defect-free ACCs is, therefore, a key factor in improving their toughness. Duchemin et al. also suggested that the mechanical properties can be affected by other factors such as temperature and humidity during preparation [67]. The degree of polymerization of the

122 Advanced Green Composites cellulose source should also be taken into account when considering the mechanical properties [61, 67, 69]. Uniaxial orientation of fillers can effectively improve the mechanical properties, and uniaxially reinforced ACCs show increased moduli and strength along the fiber direction because of the anisotropic mechanical properties of the cellulose crystal [6]. Uniaxially reinforced ACCs can be easily prepared using ramie fibers as the reinforcing phase [27, 60] because of their naturallyoriented structures. Uniaxially oriented ACC showed good Young’s modulus and strength, 20 GPa and 480 MPa, respectively, along the fiber direction. Because of the orientation and high fiber contents of the ACCs, these uniaxially reinforced ACCs have high tensile strengths and Young’s moduli in the range of 460–540 MPa and 20–30 GPa, respectively [27, 60, 81]. Mechanical drawing (stretching) of randomly reinforced ACCs can induce filler alignment (Figure  6.6). The orientation of the filler can be easily controlled to favorable levels by changing the draw ratios [64, 65]. Gindl and Keckes achieved significant improvements in the mechanical

500 nm (a)

500 nm (b)

Figure 6.6 Lengthwise sectional transmission electron micrograph images of (a) randomly and (b) uniaxially reinforced all-cellulose composites with 10% filler. Arrow indicates drawing direction. Reproduction of this figure, modified from Ref. [65], with permission from Springer (© Springer, 2016).

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properties of ACCs by drawing, achieving a maximum modulus and strength of 33 GPa and 430 MPa, respectively, along the draw direction [64]. In the ACC, the cellulose matrix, which was regenerated from LiCl/ DMAc, was also aligned after drawing, and the changes in the molecular orientation helped to improve the mechanical properties [86]. Because cellulose crystals have anisotropic diamagnetic susceptibility, magnetic field alignment is an effective technique for aligning cellulose [87]. On application of a strong magnetic field (5−20 T), the fillers align perpendicular to the magnetic field direction to some extent, and the obtained ACCs show better moduli and strengths along the aligned direction [79]. Figure  6.7 shows the mechanical properties of randomly and uniaxially reinforced ACCs. The modulus and strength increase with increasing filler content, and uniaxial reinforcement gives ACCs with higher stiffness and strength, as predicted by a finite element model [85]. The moduli and

Strength (MPa)

Modulus (GPa)

20 10

400 300 200 100

0

(a)

20

40

60

80

0

100

Filler content (%) 30

0

(b)

15 10

40

60

80

100

Uniaxially reinforced ACC Randomly reinforced ACC

Modulus (GPa)

20

20

Filler content (%)

Uniaxially reinforced ACC Randomly reinforced ACC

25 Strain at break (%)

Uniaxially reinforced ACC Randomly reinforced ACC

500

30

0

10

5 0

(c)

600

Uniaxially reinforced ACC Randomly reinforced ACC

40

1 0

20

40

60

80

Filler content (%)

10

100 (d)

100 Strength (MPa)

Figure 6.7 Mechanical properties of all-cellulose composites [27,60-65,67-70,7377,79,80]: (a) modulus, (b) strength, and (c) strain at break versus filler content, and (d) modulus versus strength.

124 Advanced Green Composites strengths of uniaxially reinforced ACCs surpass those of most nanocellulose/polymer composites, probably because of good filler/matrix binding and filler orientation. They also show a maximum strength of 480 MPa [27], which is as high as that of steel, although the density of an ACC is about 5 times lower than that of steel. ACCs show brittle behavior as a result of a significant decrease in the strain at break. Improvements in toughness, along with increases in the modulus and strength, remain a challenge in ACC preparation. As the filler content increases, the mechanical properties of randomly reinforced ACCs approach those of nanocellulose films (i.e., 100% filler fraction), which typically have Young’s moduli and strength and strain values at break of 6−15 GPa, 200−300 MPa, and 3−10%, respectively [10,88]. A uniaxially reinforced ACC with a filler content of 80% had similar properties to those of oriented nanocellulose films with Young’s modulus and tensile strength of 46 GPa and 474 MPa, respectively [89]. Generally, the mechanical properties of a material reflect the characteristics of the intermolecular interactions. In the case of ACCs, the mechanical properties are dominated by the filler content and orientation along with the strong interfacial hydrogenbonding interactions, as in conventional composites.

6.3.3 Thermal Expansion Behavior Cellulose does not melt below its thermal degradation temperature and the chains bind strongly with one another via intra- and inter-molecular interactions such as hydrogen bonds and van der Waals interactions, forming stable crystalline structures, therefore cellulose microfibrils have low coefficients of thermal expansion (CTEs) over a wide temperature range of up to ~ 200 °C [8, 9]. Yang et al. reported that the CTE of a randomly reinforced ACC film decreased with increasing filler content, and the film expanded by 6.8 ppm K−1, with only 1 wt% TOCN [74]. This CTE value is much lower than those of most plastics (>50 ppm K−1), by almost an order of magnitude [90]. Because the cellulose microfibrils show anisotropic thermal expansion, which arises from their crystalline structures, and have the lowest CTE value along the c-axis (covalently bonded direction) [8, 9], uniaxially reinforced ACCs show excellent CTEs of 0.1 ppm K−1 (80% filler), which is much lower than those of Fe and Si [27].

6.3.4 Gas Barrier Properties It has been reported that nanocellulose [91, 92] and regenerated cellulose [93] films have excellent oxygen barrier properties at low relative humidity

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(RH). The oxygen barrier properties of ACC films improve with increasing filler content [74]. Although the oxygen barrier properties become worse with increasing RH, as is the case for nanocellulose films [94], the oxygen permeabilities of ACCs at 50% RH are still better than those of commercially used high oxygen barrier polymer films such as high-density poly(ethylene), poly(ethylene terephthalate), poly(vinylidene chloride), and poly(vinyl alcohol) films [95]. Larsson et al. reported that ACC films prepared from core–shell–structured cellulose nanofibrils [96] have excellent oxygen barrier properties, even at 80% RH. Yousefi et al. reported that ACC films have complete air barrier properties (0 μm Pa−1  s−1) because of formation of fully consolidated ACC structures [62, 68].

6.3.5

Biodegradability

The inherent biodegradability of ACCs is one of their most interesting properties. Kalka et al. performed soil burial experiments to investigate the biodegradation behaviors of ACCs. They showed that the biodegradability of ACCs is better than those of other biodegradable polymers such as poly (lactic acid) (Figure 6.8) [97]. Biodegradation proceeds preferentially in the matrix phase and, therefore, the ACC biodegradability can be controlled by changing the preparation conditions. They also suggested that industrial composting is likely to be unnecessary, and this provides further motivation for developing ACCs as biodegradable composite films.

6.4 Future Prospects Because of their good mechanical properties and biodegradability, ACCs can be used as green alternatives to composite materials reinforced with glass or carbon fibers. The biocompatibility of ACCs enables their use in medical and biotechnological applications such as virus removal membranes [98] and scaffolds for tissue engineering [99]. ACC films with gas barrier properties have potential applications in food packaging [71, 100]. Recently, the concept of ACCs has been extended to cellulose/cellulose derivative composites, using cellulose acetates [101–103], carboxymethylcellulose [104], and hydroxypropylcellulose [105] as matrix phases. This could broaden the application range of ACCs. A transparent TOCN/cellulose triacetate film with high toughness, which was prepared without significantly sacrificing the low birefringence of the original cellulose triacetate, could be used in optical devices [101]. A CNC/cellulose acetate mixture can be used to prepare dry-spun microfibers [103]. This dry-spinning

126 Advanced Green Composites

ACC-u

ACC-f

Rayon-PLA

Asreceived

T170 days

T270 days

Figure 6.8 Fungicide-treated laminates are referred to as ACC-f, while the untreated laminates are referred to as ACC-u. Photographs of all-cellulose composites (30 × 30 mm2) and rayon–poly(lactic acid) composite (20 × 22 mm2) before and after soil burial testing at ambient temperature (T1) and elevated temperature (T2). Reproduction of this figure, modified from Ref. [97], with permission from Elsevier (© Elsevier, 2014).

process does not require regeneration or washing, and could enable production scale-up to a commercial scale.

6.5 Summary We have briefly described the recent progress in ACCs, which are fully renewable and bio-based composite materials. These materials are better than most commercial polymers or nanocellulose/polymer nanocomposite materials in terms of their mechanical properties because of the ideal filler/matrix binding, resulting from their identical chemistry. ACCs are prepared using only cellulose, but their mechanical properties can be finely tuned by changing the filler/matrix ratio and filler orientation. ACCs are expected to be used as alternatives in applications such as commercial

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fiber-reinforced composites, food packaging films, and optical devices. Much progress has been made in the preparation of ACCs, but the current challenge is how to achieve these excellent properties in scaled-up production processes. Moreover, because ACCs have to be dissolved in typically high-boiling-point solvents, future exploration of easier and greener ACC preparation methods is also important.

6.6 Acknowledgements This study was partly supported by Grants-in-Aid for Scientific Research from the Core Research for Evolutional Science and Technology (CREST) program of the Japan Science and Technology Agency (JST), and a Grantin-Aid for Research Activity Start-up (No.  15H06848) from the Japan Society for the Promotion of Science (JSPS).

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7 Self-Healing Green Polymers and Composites Joo Ran Kim and Anil N. Netravali* Department of Fiber Science and Apparel Design, Cornell University, Ithaca, USA

Abstract Self-healing materials such as polymers, resins, and fiber-reinforced composites have the ability to autonomously heal and restore their original properties when they are damaged through thermal, mechanical, ballistic, or other external means. During use, all materials accumulate damages such as microcracks, punctures, cuts, and scratches that result from the constant stress and strain they are subjected to. As these damages accumulate over time, they affect the properties of polymers and composites reducing their ability to function satisfactorily. This significantly cuts down their service life. Over the past 3–4 decades, researchers have developed different self-healing techniques to repair the damages as and when they occur, autonomously, and restore their properties as much as possible. The techniques developed thus far fall into three main categories: (1) capsule-based, (2) vascular, and (3) intrinsic. This chapter discusses how these three techniques work and summarizes the research done on conventional self-healing of polymers, resins, and composites, primarily using the microcapsule technique. The same self-healing techniques are now being extended, with good success, to plant derived protein and starch-based green thermoset resins as well as green composites fabricated using these resins. Self-healing green resins and composites, with increased service life, can be expected to be acceptable in mainstream applications in the future. In many cases, self-healing of green resins and composites could make themselves competitive with the conventional composites, reducing our dependence on petroleum-based materials. This would be a giant step toward cleaning up our environment. Keywords: Self-healing green resins, self-healing composites, microcapsules, cracks, fracture

*Corresponding author: [email protected] Anil Netravali (ed.) Advanced Green Composites, (135–186) © 2018 Scrivener Publishing LLC

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136 Advanced Green Composites

7.1 Introduction 7.1.1 Self-Healing Property in Materials: What is it and Why it is Needed? Self-healing materials are designed to have the ability to heal and restore their original properties when they are damaged through thermal, mechanical, ballistic, or other means [1]. They could be any materials such as polymers, metals, ceramics, or their composites [2]. Polymers and their composites are commonly used in wide ranging fields such as aerospace, ground transportation, civil engineering, electronics, coatings, adhesives, and sport goods. However, over their working life, these materials continue to accumulate damages from constant stress and strain such as cracking, puncture, cut, corrosion, delamination, and fiber debonding. A schematic of the damages that can occur in polymer composites is shown in Figure 7.1 [3]. As these damages accumulate, they significantly affect their service life. Polymers and composites, in particular, are known to undergo losses in their mechanical properties under constant strain and stress experienced during use. This property reduction is primarily due to the formation of microcracks that continue to propagate when exposed to common elements of the atmosphere such as moisture and oxygen. Propagating cracks accelerate the property degradation and ultimately render the materials useless for their intended application [4]. It has been reported that the total annual estimated direct cost from this type of microcrack-related damage, in many nations, can be between 2% and 5% of the nation’s gross domestic

b

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Deep cut in coating Corrosion in protected metal

i

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Figure 7.1 Schematic of the damage modes in polymer composites [3].

Self-Healing Green Polymers and Composites 137 product (GDP) [5]. In order to overcome microcrack-related limitations, many types of self-healing materials have been developed and studied. These materials offer safer and longer service life for the products and have become a commercial reality. Self-healing concept was first introduced by White et al. in 2001 [6], and it was defined as “a material that can heal the damages autonomously according to a healing response,” that is, a material that has the ability to repair itself [7]. Self-healing mechanism has been classified into two distinct types: extrinsic and intrinsic. Extrinsic self-healing system is achieved with the help of embedded healing agents; whereas, intrinsic self-healing system is achieved by the polymer itself [8]. Intrinsic self-healing can consist of reversible reactions involved in reconnecting intra- and/or intermolecular linkages such as radical thiuram disulfide units [9], disulfide links [10, 11], C–C bonds [12], hydrogen bonding [13], - stacking [14], covalent acylhydrazone bonds [15], and Diels–Alder reactions [16] triggered by light and/or heating. Extrinsic self-healing system is achieved through irreversible reactions. In such cases, the healing or cross-linking agents can be stored in capsules [17, 18] or hollow fibers until needed, that is, until the microcracks are formed [19, 20]. Several concepts of extrinsic self-healing system have been developed to provide self-healing property to the material, such as microcapsule-based systems, microvascular systems, and nanoreservoirs [21]. Microcapsule-based self-healing has become the preferred method because of facile and economical capsule production process as well as the ease of incorporating them into many polymers, composites, and biomaterials [3]. The first part of this chapter focuses on self-healing principles and methods using microencapsulated capsules that repair the damage as soon as the microscopic cracks are formed, as at this level the damage often goes unnoticed. Quick damage repair enables the material to restore most of its mechanical properties. The latter part of the chapter is devoted to self-healing of green resins and green composites.

7.2

Types of Self-Healing Approaches Used in Thermoset Polymers

In order to overcome the limitations of polymers and composites for their final applications, many types of self-healing materials have been developed that can increase the useful life of the products [4]. Three methods, for example, microcapsule [22, 23], microvascular [24], and intrinsic systems, have been successful in providing self-healing function to a variety of polymers. These methods are depicted in Figure 7.2 [3, 21, 25].

138 Advanced Green Composites Capsule based

Vascular

Intrinsic

(a)

(b)

(c)

Figure 7.2 Three types of commonly used self-healing approaches: (a) capsule-based, (b) vascular, and (c) intrinsic methods [3].

7.2.1 Microcapsule-Based Self-Healing System Among the three systems, the microcapsule-based system has been the most popular method used to accomplish self-healing of a variety of materials [26, 27]. In this system, the healing agent is stored in discrete microcapsules that are dispersed through the resin. When the damage such as microcracks form, they trigger the rupture of the microcapsules that are present in the path of the microcracks releasing the healing agent. The released healing agent reacts with hardener or cross-linking agent present in the resin/polymer. The healing agent and the cross-linking reactions with the help of the polymer restore the original mechanical properties, shape, or appearance of the polymer, at least partially, if not fully [3]. The encapsulation in self-healing system have been introduced using interfacial, in situ, coacervation, sol–gel, or physical methods such as spray drying and solvent evaporation on the basis of the mechanism of wall formation [28]. In situ and interfacial methods have been used for encapsulating epoxy healing agent using shell materials consisting of urea-formaldehyde (UF) [29], melamine formaldehyde (MF) [30, 31], melamine-ureaformaldehyde (MUF) [32], polyurethane (PU) [33], or acrylates [34]. Shell materials such as UF, MF/MUF, and PU have been shown to provide great durability as they are able to survive the processing conditions and manufacturing processes and can be prepared in many desirable sizes [29–38]. White et al. focused on self-healing epoxy resins using dicyclopentadiene (DCPD)-loaded poly(urea-formaldehyde) (PUF) microcapsules while dispersing Grubbs’ catalyst in the resin [35]. When the microcrack damage is initiated, the DCPD-PUF microcapsules are punctured, releasing DCPD into the epoxy resin. On contact with the first generation Grubbs’ catalyst that is present in epoxy resin, DCPD undergoes ring opening metathesis polymerization (ROMP) with DCPD and heals the microcracks as shown

Self-Healing Green Polymers and Composites 139 CI CI DCPD monomer

PCy3 Ru

Ph H

PCy3

Grubbs’ catalyst

Cross-linked polymer network

Figure 7.3 Chemical self-healing reaction of DCPD and Grubbs’ catalyst resulting in cross-linking network in the resin [5].

in Figure 7.3. Their results showed excellent fracture strength recovery of about 67% in case of virgin epoxy resin [35]. Polydimethylsiloxane (PDMS)-loaded PUF microcapsules have also been reported to withstand the harsh processing conditions in comparison to other shell materials because PUF can withstand much higher temperatures than UF, MF, and PU [36]. Other encapsulation methods in self-healing system have been also studied using solvent evaporation method [37], Pickering stabilization [38], inverse Pickering stabilization [39], and multiple emulsions [40]. In the “single capsule” system, microcapsules containing healing agent are dispersed in the resin that already contains the catalyst needed for the reaction [41–60]. In this system, the healing agents released from capsules flow into the microcracks and bridge them by chemical and physical interactions with the catalyst in the resin. In comparison, the “double capsule” system involves preparing two types of capsules that encapsulate the healing agent and catalyst individually and dispersing them into the resin. When the capsules are ruptured, the released healing agent and catalyst come in contact with each other as well as the resin and polymerize/crosslink within the voids or microcracks in the resin and healing it. Table 7.1 summarizes currently used self-healing microencapsulation systems and provides details such as single, double, or multi-layered microcapsule systems, healing agent used, shell material, the type of resin, catalyst, the healing efficiency obtained, and the healing conditions.

7.2.1.1 Microencapsulation Techniques Microencapsulation may be defined as a process to package or cover particles derived from ground solids, droplets of liquids, or gaseous materials with protective membranes that are commonly called as shells or coatings. Microencapsulation techniques are categorized commonly by physical and chemical methods. Chemical methods such as in situ interfacial and sol–gel polymerization require that two reactants meet at the interface of emulsion and react rapidly by polycondensation. The physical

PUF

Poly (phenylene oxide

PUF

PMF PMF

PUF

Epoxy

Epoxy

Glycidyl methacrylate (GMA)

Dibutylphthalate

PMMA

Epoxy PMMA PMMA composites

Cyanate ester

4,4 -bismaleimido- diphenylmethane (BDM)/ diallylbisphenol A

Epoxy

Shell material Resin

Epoxy

Single capsule system

Healing agent

Excess amine (DETA)

Catalyst

6 wt%/100%/RT for 72 h

10 wt%/75%/RT for 72 h 15 wt%/100%/25 °C 21 h 10 wt%/120%/25 °C for 72 h

[48]

[45] [46] [47]

[44]

[43]

10 wt%/79%/220 °C for 5 h 5 wt%/85%/1 h at 220 °C

[41] [42]

References

15 wt%/82%/*RT 24 h 20 wt%/100%/RT 24 h

Healing agent concentration/healing efficiency/healing condition

Table 7.1 Encapsulation systems used for self-healing materials, the healing agents and shell material, a catalyst, healing conditions, and healing efficiency in single or double capsule-based system.

140 Advanced Green Composites

PUF PUF

PUF

Linseed oil Tung oil

Methacryloxypropylterminated polydimethylsiloxane

PUF

MUF

PUF

PUF

Dicyclopentadiene (DCPD)

5-ethylidene2-norbornene (ENB)

DCPD and ENB

DCPD

Single with disperse catalyst

Polyurethane

2-octylcyanoacrylate

PEO

PEO

Epoxy resin polymerized

Epoxy

Poly(dimethylsiloxane-comethylphenylsiloxane) +tetraethyl orthosilicate

Paint film Epoxy coating

Acrylic bone PMMA

5% Grubbs’ catalyst

5% Grubbs’ catalyst

Diethylene tetramine (DETA) curing agent

20 wt%, 2.5 wt% Catalyst/213%/ 30 °C for 24 h

–

–

15 wt%/63%/RT for 24 h

(Continued)

[56]

[55]

[54]

[53]

[22]

12.5 wt%/Recovered by observation/ under sunlight for 4 h

Denzoin isobutyl ether (photo initiator)

2.5% Grubbs’ catalyst

[51] [52]

15 wt%/Recovered by observation/RT for 72 h 12%/Recovered by observation/10 days

[49] [50]

Tung oil

6.8 wt%/70%/RT for 72 h

Self-Healing Green Polymers and Composites 141

PUF

PMF

PUF

Epoxy+ CuBr2

Iodonium bis(4-methylphenyl) hexafluorophosphate + GMA

Mono- or di-Norbornenedicarboximide

Epoxy monomer

PUF

Epoxy

Double capsules system (catalyst in capsules)

Epoxy

PS composites

Epoxy composite Epoxy composites

Shell material Resin

Healing agent

Table 7.1 Cont.

Aliphatic polyamine

Grubbs’ catalyst ruthenium initiator (5 MOL%) or RuCl2

4% NaBH4

2% CuBr2(2-MeIm)4

Catalyst

40 wt%/77%/RT for 24 h 10.5 wt%/91%/RT for 48 h 15%/62%/50 °C for 24 h

10 wt%/33%/RT for 24 h

10 wt%/100%/25 °C for 24 h

30 wt%/93%/140 °C for 30 min 10 wt%/94%/140 °C for 30 min with compression

Healing agent concentration/healing efficiency/healing condition

[61] [61] [62]

[60]

[59]

[57] [58]

References

142 Advanced Green Composites

Epoxy vinyl ester matrix

PU

PUF

PUF

Hydroxyl endfunctionalized polydimethylsiloxane (HOPDMS) and polydiethoxysiloxane (PDES)

Poly(dimethyl siloxane) (PDMS)

Methylmethacrylate

Epoxy vinyl ester resin

PDMS

PVC

PMF

GMA and polythiol pentaerythritol tetrakis (3-mercaptopropionate)

Epoxy resin Epoxy composite

PUF

Epoxy

Tertiary amine and benzoyl peroxide (BPO)

Platinum

Dimethyl dineodecanoatetin (DMDNT)

2,4,6-tris (dimethylaminomethyl) phenol catalyst

(C2H5)2O·BF3 catalyst

15 wt%/80%/RT for 48 h

(Continued)

[68]

[67]

[66]

12 wt%, 3 wt% catalyst/recovered by observation/RT for 24 h

20 wt%, 5 wt% catalyst/100%/RT for 24 h

[65]

[63] [64]

56 wt%/90%/25 °C for 3 h

5 wt%, 1 wt% catalyst/81%/20 °C for 30 min 5 wt% and 1 wt% catalyst/76%/20 °C for 30 min

Self-Healing Green Polymers and Composites 143

PUF

PMF

Epoxy/pentaerythritol tetrakis(3-mercaptopropionate)

Azide-modified poly(isobutylene)

*RT is the room temperature

GMA and PMMA

PMF

PMMA

Epoxy

Multi-layers capsules

PUF

Styrene

Polystyrene resin

Poly(isobutylene)

Epoxy

Epoxy

Epoxy

Shell material Resin

Healing agent

Table 7.1 Cont.

N,N,N,N,N pentamethyldiethylene triamine and CuCl

Trivalent alkynes and CuIBr(PPh3)3

Tertiary amine

Polyetheramine

Benzoyl benzenecarboperoxoate

Catalyst

[72]

[71]

10 wt%, 5 wt% catalyst/91%/5 days

15 wt%/100%/RT for 48 h

[30]

[70]

[69]

References

5 wt%/104.5%/20 °C for 24 h

7.5 wt%, 7.5 wt% catalyst/84.5%/RT for 24 h

15 wt%, 3 wt% catalyst/65%/RT for 24 h

Healing agent concentration/healing efficiency/healing condition

144 Advanced Green Composites

Self-Healing Green Polymers and Composites 145 methods such as solvent evaporation, centrifugal extrusion, polymer extrusion, coating, and spray drying have been used to fabricate microcapsules without the occurrence of chemical reactions at the interface. These methods apply coatings to the surface of the substrate [73]. The most widely used methods in self-healing systems are described in more details below. 7.2.1.1.1 In Situ Polymerization In situ polymerization is a chemical encapsulation technique, where direct polymerization of a single monomer is carried out at the core/shell interface [28]. Typically, the microencapsulation processes can start either directly through cross-linking reaction between amine and aldehyde monomers or by changing pH or temperature, forming aminoaldehyde microcapsule wall from the precondensates. When aminoaldehyde precondensates are used for in situ polymerization microencapsulation, the process itself results in the formation of a shell. From the summary presented in Table 7.1, it can be seen that MF, UF, PUF, or poly(melamine-formaldehyde) (PMF) have been good candidates as shell materials for in situ polymerization. This method is similar to the interfacial polymerization method explained below. The distinctive feature of in situ polymerization, however, is that no reactants are included in the core material. The polymerization occurs in the continuous phase, rather than on two sides of the interface between the continuous phase and the released core material [28]. Examples of this method include PUF and PMF encapsulation systems [74]. Most microcapsules developed using in situ polymerization method, have been in oil-in-water emulsions. The produced microcapsules display spherical shapes with smooth surfaces without any pores [75]. Using the in situ method, PUF has been used as a shell to encapsulate DCPD. However, it showed brittle behavior and poor sealing properties and also low weather resistance [74]. In another example of in situ encapsulation method, PMF was used to encapsulate diglycidyl tetrahydro-o-phthalate, as a healing agent, which showed better sealing properties than PUF [5]. However, PMF is expensive and, hence, is difficult to be used on commercial scale. On the other hand, poly(melamine-ureaformaldehyde) (PMUF) has received significant attention in the past couple of decades because of its cost-effectiveness [32]. Polyphenylene oxide (PPO) has also been developed to encapsulate diglycidyl ether of bisphenol A (DGEBA) using the in situ method [76]. A significant advantage of using PPO is that it can withstand temperatures of up to 258 °C; whereas, PU, PMF, and PMUF soften above 200 °C [8]. The double layer microcapsules that contain epoxy and hardener (diaminodiphenyl sulfone) have also been developed in single capsule system via in situ polymerization [77].

146 Advanced Green Composites 7.2.1.1.2 Interfacial Polymerization Interfacial polymerization (IFP) for microencapsulation is characterized by wall formation via rapid polymerization of monomers at the surface of dispersed core material [78]. In this case, two reactants used in the polycondensation reaction meet at the interface and react rapidly. Typically, a multifunctional monomer is dissolved in the core material, and the solution is dispersed in an aqueous phase. A reactant that will react with the monomer is then added to the aqueous phase allowing the polymerization to occur quickly at the interface, forming the microcapsule walls [28]. IFP can be useful for preparing larger microcapsules, compared to other methods. Most commercial IFP processes produce capsules in 20–30 μm range for antimicrobial applications as well as smaller ones, 3–6 μm, for carbonless paper ink applications [37, 79]. The IFP process has many advantages such as fast encapsulation and high encapsulation efficiency. However, it is more likely to contain unreacted shell monomers. As an example, 2-octylcyanoacrylate (OCA) has been encapsulated as a healing agent using IFP process [80]. In addition, epoxy-loaded ethylenediamine microcapsules have been developed by IFP process. Epoxidized linseed oil-loaded PU microcapsules have also been produced by IFP using two solvents, methylene diphenyl dissocyanate and dendritic polyamidoamine [81]. 7.2.1.1.3 Sol–Gel Technique The sol–gel technique allows encapsulation of inorganic or lipophilic materials, as core, inside silicon dioxide shell materials. In this process, once the inorganic molecular precursors are dissolved in solvents with low viscosity, shell reactant is added to the solution. The mixture gradually becomes a gel [8]. Sol–gel microencapsulation then occurs in the emulsion state. Emulsions generally consist of two immiscible phases that are stabilized using a surfactant. Emulsions are commonly prepared by dispersing one liquid into another in the form of fine droplets. Commonly used emulsion systems are oil-in-water (o/w) or water-in-oil (w/o) types, both of which can be stabilized and have been found to be useful in the preparation of microcapsules [82]. One example is an o/w emulsion with a core material solubilized in the silicon phases such as tetraethoxysilane or tetramethoxysilane [82]. In this case, hydrolysis of silicon droplets and condensation of hydrolyzed silicon to silica occur at the interface. These reactions occurring at the interface lead to the formation of shells made of silica. The conditions of the process can be easily manipulated with the help of several variables such as surfactants, nature of silicon precursors, temperature, time, and mixing conditions [82].

Self-Healing Green Polymers and Composites 147 Silica coating

Self-healing catalyst particles

517 nm

1.8 um

Self-healing monomer capsules

Figure 7.4 DCPD-loaded silica microcapsules produced using sol–gel encapsulation [83, 84].

Microcapsules using silica as shell material in self-healing systems have been investigated using the sol–gel encapsulation method [83, 84]. In this self-healing system, DCPD-loaded microcapsules produced using sol–gel encapsulation and dispersed Grubbs’ catalyst present in the epoxy resin displayed good protection of the core material and reduced their aggregation as shown in Figure 7.4. The self-healing in this study was seen to be adequate. 7.2.1.1.4 Solvent Evaporation Technique Solvent evaporation technique is a physical encapsulation method that involves emulsifying volatile organic solvent in water that is followed by solvent removal. Initially, the shell polymer solution is prepared by dissolving the polymer in a volatile solvent, which is immiscible with the core material. Depending on the hydrophobicity or hydrophilicity of the core material, it is dispersed in the shell polymer solution and stirred continuously until the solvent is partitioned into the aqueous phase and evaporated. The shell material shrinks around the core material and results in hardened microspheres through drying of the shell materials [85]. Compared to other techniques, solvent evaporation technique has many advantages such as being economical, facile, and diverse in the selection of core and shell materials [64]. In solvent evaporation methods, w/o emulsification process is the most widely used process to encapsulate hydrophilic core materials. However, it contributes to large average diameters of microcapsules and low encapsulation efficiency [86–88]. This method is commonly

148 Advanced Green Composites used to encapsulate water-soluble proteins, vaccines, and drugs in many pharmaceuticals [89, 90]. Furthermore, Ogawa et al. developed a waterin-oil-in-water (w/o/w) double emulsion solvent evaporation method to increase core loading of microcapsules using poly(d,l-lactide-co-glycolide) (PLGA) as the shell [91]. This method showed high protein loading, and encapsulation efficiency as compared to w/o solvent evaporation method [89]. In the w/o/w method, the water-soluble protein in water is added to an organic phase such as dichloromethane (DCM) containing PLGA and then stirred with a high speed to produce the first w/o emulsion. The w/o emulsion is then put into a large amount of water containing an emulsifier such as poly(vinyl alcohol) (PVA) or poly(vinyl pyrrolidone) (PVP) to produce the w/o/w emulsion using high speed stirring. The emulsion is then subjected to solvent removal either by evaporation or extraction process. In the former case, the emulsion is maintained at reduced pressure or at atmospheric pressure and stirred to enable DCM to evaporate. The solid microcapsules obtained are then washed and collected by filtration and centrifugation [92]. By using the w/o/w solvent evaporation method, double microcapsule system, DGEBA epoxy, and polyetheramine (hardener) with PLGA shell have been developed. The core loading was over 50 wt% and 20 wt% of epoxy and hardener, respectively [70]. Various shell materials using solvent evaporation technique are commercially available and include poly(L-lactide), PMMA, PLGA, poly(vinyl formal), poly(vinyl acetate), PPO, and poly(vinyl cinnamate) [8]. In another study involving solvent evaporation technique, DCPDloaded microcapsules within silica oxide shell were produced using an o/w solvent evaporation method for self-healing applications [83]. In this method, the core material (healing agent) should be hydrophobic so that the water can be removed after production of microcapsules.

7.2.1.2 Microcapsule Systems for Self-Healing 7.2.1.2.1 Single Microcapsule System In a single capsule system used for self-healing materials, microcapsules containing the healing agent or catalyst are uniformly dispersed within the resin in composites [3]. The component not in the microcapsules, needed for reacting or cross-linking, may be inside the composite, that is, resin. Liu et al. studied the microcapsule–catalyst system using DCDP as the healing agent [93]. DCPD has been reported to react fast and needing minimum amount of catalyst to produce high Tg without a distinct melting point [55]. Moll et al. developed a self-sealing composite by adding DCPD-loaded microcapsules and Grubbs’ catalyst into a glass

Self-Healing Green Polymers and Composites 149

Epoxy matrix

GMA HO OH N GMA

GMA

GMA

Epoxy matrix

N

N

GMA

OH

O

OH

N

N

HO

GM A

HO

O

GMA

Epoxy matrix

OH GMA

O

HO

Epoxy matrix

fiber-reinforced epoxy composite [94]. Rong et al. developed self-healing epoxy resin containing epoxy-loaded microcapsules and imidazole as catalyst [95]. Glycidyl methacrylate (GMA)-loaded microcapsules in epoxy resins have also been developed for self-healing applications [45]. In this system, the released GMA was able to cross-link microcracks even at RT through hydrogen as well as covalent bonds and restore fracture toughness with high healing efficiency as shown in Figure 7.5 [45]. GMA contains epoxy groups and C=C bonds inside microcapsules, so nucleophilic addition and ring-opening reactions between GMA and the amine curing agent present in epoxy resins occur as microcracks propagate, reconnecting the fracture surfaces [45]. PMMA resins containing GMA-loaded microcapsules, as the healant, showed that the healing efficiency (impact strength) increased rapidly with increasing healing time. These studies resulted in 89% and up to 100% recovery after 12 h and 21 h healing times, respectively [46]. Another study using GMA-loaded microcapsules in PMMA composites also showed high healing efficiency of almost 100% after 72 h of healing at RT by reversible addition–fragmentation chain transfer (RAFT) polymerization [47]. In yet another study, both 10 wt% 4,4 -bismaleimido-diphenylmethane (BDM) and diallylbisphenol A(BA) filled poly(phenylene oxide) microcapsules (PPOMCs) in epoxy resins showed healing efficiency of 79% after heat treatment at 220 °C for 5 h [43]. It has been proven that the PPOMCs can be controlled by both heat and force to release the healing agent, which polymerizes in the presence of various catalytic and reactive groups such as amine and hydroxyl groups to cross-link the microcracks in

N GMA HO

O

GMA

GMA

(a)

OH

(b)

Figure 7.5 Microcrack healing in epoxy composite containing GMA-loaded microcapsules. Hydroxyl and ternary amine groups can cross-link on the fracture surfaces: (a) distribution of the released GMA in the damaged region and (b) hydrogen bonds formed between the microcrack surfaces [45].

150 Advanced Green Composites the resin. This bonding of microcrack surfaces allows mechanical property recovery of resins and increases their service time [43]. Epoxy resin-loaded microcapsules in cyanate ester resin have also been studied [44]. The self-healing ability of cyanate ester resin containing just 5 wt% microcapsules had 85% healing efficiency compared to their original fracture toughness after 1 h recovery period at 220 °C [44]. Linseed oilloaded UF microcapsules have been studied in epoxy paint films. Cracks in paint films were successfully healed as healing agent was released from broken UF microcapsules [51]. Figure 7.6 shows the recovery or healing of a typical scratch in epoxy paint films that contain linseed oil-loaded UF microcapsules, in just 90 s, under optical microscopy. Similar self-healing coating study has been done using Tung oil-loaded UF microcapsules [52]. Released Tung oil from ruptured microcapsules in epoxy resin film showed self-healed scratches after 72 h, at RT, as shown

Initial

15 seconds

30 seconds

45 seconds

60 seconds

90 seconds

Figure 7.6 Optical microscopy images of healing of a typical scratch in epoxy paint films containing linseed oil-loaded UF microcapsules for 90 s [51].

Self-Healing Green Polymers and Composites 151 in Figure 7.7 [52]. Figure 7.7(a) and (c) shows separated or heavily carved scratches at 0  h. These gapes and deep scratches, however, were filled in 72 h due to released Tung-oil, used as the healing agent, as shown in Figure 7.7(b) and (d). Light-activated self-healing polymer has also been developed using single microcapsule-based system. Both methacryloxypropyl-terminated polydimethylsiloxane (MAT-PDMS) as the healing agent and benzoin isobutyl ether (BIE) photoinitiator with methyl cinnamate as the fluorescent indicator were encapsulated in UF-based microcapsules using sol–gel method. The microcapsules were dispersed in tetraethyl orthosilicate (TEOS)-based resins. In order to observe the healing phenomenon, healed cracks were observed under fluorescence microscopy in TEOS resin [22]. The microcapsules were seen to self-heal by the release of the healing agent from microcapsules promoted (prompted) by sunlight, as shown in Figure 7.8. The solvent-promoted self-healing was developed by Caruso et al. [97,98]. In this study, 20 wt% epoxy in chlorobenzene within UF microcapsule shells was prepared using in situ polymerization. The epoxy resin containing 15 wt% of microcapsules showed 82% healing efficiency [41]. This research has been further expanded using two green solvents, phenyl

(a)

(b)

400 m

200 m

(c)

(d)

Figure 7.7 Healed epoxy film containing Tung oil-loaded UF microcapsules: (a) artificial scratches before healing, (b) healed scratches after 72 h, (c) SEM image of a scratch before healing, and (d) SEM image of a scratch after 72 h healing [96].

152 Advanced Green Composites

500 m (a)

500 m (b)

Figure 7.8 Self-healing of resins prompted by sunlight: (a) after healing, and (b) before healing of the microcrack under fluorescence microscopy [22].

acetate, and ethyl phenylacetate, which have lower toxicity than chloroform. The incorporation of an epoxy monomer in these two solvents to form filled microcapsules resulted in improved healing efficiency of up to 100% in epoxy resins [42]. Most commonly used polymer, PMMA, as bone cement, is a spacefilling resin that initiates mechanical interlocks between the stem of the implant and the surrounding boney tissue. Bone cement consists of two components: low molecular weight PMMA polymer and an initiator such as benzoyl peroxide or liquid methyl methacrylate monomer [99,100]. Mixing PMMA with the initiator produces the space-filling resin that is used in implants [99,100]. The most common problem in biomaterial applications is short service time due to high friction of the bone joints [68]. As a result, self-healing systems using a single microcapsule have been used in biomaterials such as bone cement and joint applications [68]. Dibutylphthalate (DBP)-loaded UF microcapsules have been tested by Jackson et al. on a PMMA resin for bone cement [48]. After 72 h of healing, the crack plane and mechanical healing were observed [48]. Self-healing

Self-Healing Green Polymers and Composites 153

Load, P(N)

80 Brittle virgin fracture

60 40

Non-linear, ductile heated fracture

20 0

0

(a)

1.5 0.5 1 Displacement, (mm)

2

Load, P(N)

80 Brittle virgin fracture

60 40 20 0

(b)

Brittle healed fracture

0

0.5 1 1.5 Displacement, (mm)

2

Figure 7.9 Typical load–displacement curves of virgin and healed epoxy resins with (a) 5 wt% DCDP-loaded microcapsules containing 5 wt% Grubbs’ catalyst and (b) 2.5 wt% Grubbs’ catalyst [53].

acrylic bone cement has also been studied using 2-octylcyanoacrylate (OCA)-loaded PU microcapsules. In this case, after 24  h of healing, the acrylic bone cement showed good recovery or healing of cracks [49]. Recently, in a set of studies, recovery of electrical conductivity using the self-healing mechanism was characterized using carbon nanotubeloaded microcapsules. Microcapsules containing suspensions of polymerstabilized carbon nanotubes and/or graphene flakes were successfully used for autonomous restoration of electrical conductivity in epoxy and polyurethane resins [101,102]. DCPD-loaded UF microcapsules and Grubbs’ catalyst were studied in the epoxy resin for self-healing [53]. Figure 7.9(a) shows typical load– displacement curves of virgin and healed epoxy resins with 5 wt% DCDPloaded microcapsules containing 5 wt% Grubbs’ catalyst and Figure 7.9(b) shows virgin and healed epoxy resins having 2.5 wt% Grubbs’ catalyst. This composition resulted in up to 63% healing efficiency after 24  h healing time at RT [53]. In another study, two different healing agents, 5-ethylidene-2-norbornene (ENB) and ENB with 10 wt% norbornene-based cross-linking agent

154 Advanced Green Composites were developed using MUF shell by in situ polymerization [54]. The MUF microcapsules, prepared in this study, demonstrated greater thermal stability of up to 300 °C, compared to UF microcapsules [54]. Mixture of the mono- and di-functional norbornene dicarboximide monomers as the healing agent was successfully encapsulated in UF by majchrzak et al. [60]. Microcapsules at just 10 wt% and Grubbs’ ruthenium initiator at 5 mol% in epoxy resin showed healing efficiency of up to 33% with the expectation that higher amount of microcapsules could increase the healing efficiency further [60]. The effects of epoxy-loaded UF microcapsules and the complex of CuBr2 and 2-methylimidazole have been studied with extra hardener added to the epoxy resin. At a concentration of 10 wt% microcapsules and 2 wt% hardener, the epoxy resin was able to completely recover its original fracture toughness after curing at 130 °C for 1  h [57]. Another study using CuBr2(2-MeIm)4 as latent hardener in epoxy composites reinforced with woven glass fabrics showed 70% healing efficiency. SEM image of fracture surface of self-healing is shown in Figure  7.10 [103]. In these composites, 30 wt% epoxy-loaded UF microcapsules and 2 wt% CuBr2(2-MeIm)4 latent hardener were dispersed in the resin, and high healing efficiency was obtained after 30 min at 140 °C [103]. In a recent experiment, a mixture of iodonium bis(4-methylphenyl) hexafluorophosphate (IBH)/GMA-loaded PMF microcapsules (10 wt%) and NaBH4 catalyst (4 wt%) when embedded in the polystyrene composite resulted in 100% healing efficiency at RT for 24 h healing time [59].

Figure 7.10 SEM image of fracture surface of self-healing epoxy composite reinforced with woven glass fabrics [103].

Self-Healing Green Polymers and Composites 155 7.2.1.2.2 Dual Microcapsule System In the dual microcapsule system, both the healing agent and the crosslinking agent, either a catalyst or a hardener, are encapsulated individually. The dual microcapsule system in thermoset epoxy resin is shown using epoxy and amine hardener (curing agent) microcapsules in Figure  7.11 [61]. One microcapsule type contained a modified aliphatic polyamine hardener while the second microcapsule type contained epoxy monomer (the healing agent) at the ratio of 4:6 (w/w) for amine:epoxy microcapsule, resulting in healing efficiency of up to 77% [61]. Keller et al. demonstrated that dual capsule self-healing system using 15 wt% PDMS and 5 wt% platinum catalyst-loaded UF microcapsules in PDMS resin resulted in up to 100% healing efficiency [104]. Cho et al. were able to successfully extend the dual capsule self-healing system used for PDMS resin to epoxy coatings [105]. The epoxy coating containing PDMSloaded UF microcapsules and dimethyldineodecanoate tin (DMDNT) catalyst-loaded UF microcapsules covered the scratches at RT, within 24 h. Xiao et al. also demonstrated that epoxy resins containing two types of microcapsules, boron trifluoride diethyl etherate-loaded microcapsules and epoxy resin-loaded microcapsules, resulted in 76% healing efficiency after 30 min at 20 °C (RT) [106]. Epoxy-loaded UF microcapsules (5 wt%) as the healing agent and 1 wt% of (C2H5)2O·BF3-loaded UF microcapsules in epoxy resins containing triethylenetetraamine (TETA) hardener exhibited over 80% healing efficiency [45]. In another study, two microcapsules were embedded for self-healing of polyvinyl chloride (PVC) resin [65]. One of them encapsulated a mixture of healing agents, GMA and polythiol pentaerythritol tetrakis (3-mercaptopropionate) (PETMP), within PMF shell and

Epoxy matrix

Matrix capsule

Amine capsule 150 m

Figure 7.11 The dual microcapsules system contained both epoxy and polyamine microcapsules in the epoxy resin [61].

156 Advanced Green Composites the other encapsulated a catalyst, 2, 4, 6-tris(dimethyl aminomethyl) phenol (DMP-30) within PMF shell [65]. GMA and PETMP-loaded microcapsules and DMP-30-loaded microcapsules were added to PVC resins at the mass ratio of 1:1.56 for a total of 20wt% of PVC. Healing efficiency was up to 90% at RT after 3  h of healing [65,107]. In a study of epoxy-based composites, epoxy-loaded microcapsules, and boron trifluoride diethyl etherate, (C2H5)2O·BF3-loaded microcapsules were incorporated into sisal fiber/epoxy composites. As a result of the epoxy–BF3 cationic chain polymerization, the epoxy composites exhibited 76% healing efficiency in impact strength after 30 min at RT [64]. In another research involving dual microcapsule technique, hydroxyl end-functionalized polydimethylsiloxane (HOPDMS) and polydiethoxysiloxane (PDES)-loaded PU microcapsules as a healing agent and DMDNT catalyst solution-loaded PU microcapsules were synthesized [66]. Epoxy vinyl ester coating containing 12 wt% healing agent, 3 wt% catalyst capsules, and 3 wt% methylacryloxy propyl triethoxy silane adhesion promoter was observed to almost fully cover the scratches at RT after 24 h as shown in Figure 7.12 [66]. Figures 7.12(a) and (c) show optical microscopy and SEM images before healing, and images in Figure 7.12(b) and (d) show optical microscopy and SEM images after 24 h healing. Self-healing PDMS elastomer system has also been developed using dual microcapsule system containing platinum-catalyzed hydrosilylation-loaded

20 mm (a)

20 mm (b)

50 m (c)

50 m (d)

Figure 7.12 Self-healing epoxy vinyl ester coating after 24 h healing at RT: (a) optical microscopy image of scratches before healing, (b) optical microscopy image of scratches after healing, (c) SEM image of a scratch before healing, and (d) SEM image of a scratch after 24 h healing [66].

Self-Healing Green Polymers and Composites 157 PUF microcapsules and vinyl-terminated PDMS-loaded PUF microcapsules [67]. This resulted in 100% healing efficiency [67]. Styrene-loaded and benzoyl benzenecarboperoxoate-loaded microcapsules prepared, separately, by in situ polymerization were incorporated into epoxy resins in yet another study on epoxy resin. After adding 15 wt% styrene-loaded capsules and 3 wt% benzoyl peroxide-loaded microcapsules, 65% healing efficiency was achieved after 24 h healing at RT [69]. There are many studies on epoxy self-healing. Dual capsule system using epoxy-loaded microcapsules as the healing agent and polyetheramine-loaded microcapsules as hardener was developed in another study [70]. The epoxy resin containing 5 wt% epoxy-loaded microcapsules and 7.5 wt% polyetheramine-loaded microcapsules showed about 44% healing efficiency. After increasing the epoxy-loaded microcapsule loading to 7.5 wt%, the healing efficiency increased to over 84% at RT after 24 h [70]. Self-healing epoxy using mixture of pentaerythritol tetrakis(3-mercaptopropionate) and tertiary amine-loaded microcapsules as the hardener has also been developed as well [30]. In this system, epoxy resin containing just 1 wt% microcapsule loading and 2.5 wt% hardener showed about 44% healing efficiency and after increasing microcapsule loading to 5 wt%, the healing efficiency increased to 103%, at 20 °C after 24 h. The results of this study are shown in Figure 7.13 as load–displacement curves of healed and virgin specimens [30].

100

Load (N)

75

Healed neat epoxy (the healing agent was Stick-slip failure pre-mixed and then injected into the cracked planes)

50 Virgin self-healing specimen

25

Healed self-healing specimen (healing efficiency = 103.3%)

Neat epoxy

0 0.00

0.25

0.50

0.75

1.00

Displacement (mm)

Figure 7.13 Typical load–displacement curves of self-healing composite contains 5 wt% epoxy-loaded microcapsules and 2.5 wt% the mixture of pentaerythritol tetrakis(3mercaptopropionate) and tertiary amine-loaded microcapsules as the hardener resulted in 103.3% healing efficiency after 24 h at 20 °C [30].

158 Advanced Green Composites Protective shell Catalyst system Isolating shell Healant

Figure 7.14 Scheme of multi-layer single capsules containing healing agent and catalysts [72].

7.2.1.2.3 Multi-Layer Microcapsule System A novel self-healing system that encapsulates a mixture of all needed components such as healing agent, cross-linking agent, and catalyst in multiple layers within a single capsule was developed by Zhu et al. using PMF for shell [72]. Scheme of multi-layer microcapsules containing healing agent and catalysts is shown in Figure 7.14. In this study, with a mixture of GMA, PMMA-Br, and CuCl/N,N,N,N,N-pentamethyldi-ethylenetriamine) in polystyrene resin exhibited 100% healing efficiency after 48 h of healing at RT [72]. This study and the study mentioned in the last paragraph by Yuan et al. [30] clearly indicate that it is possible to achieve 100 self-healing efficiency, that is, no loss in properties, in epoxy resins using microcapsules.

7.2.2 Vascular Self-Healing System Vascular (microvascular) self-healing system, as shown earlier in Figure 7.2, contains the healing agent in a network of capillaries or hollow channels, which may be interconnected until damage triggers the selfhealing action [3]. Such a system is more suitable for large damage volume [3]. The microcapsule system is incapable of healing the microcracks at the same site because of the limited healing agent availability. Adding large amounts of microcapsules to overcome the deficiency, however, can reduce the properties of the resins significantly. For this reason, vascular self-healing systems have been developed that can compensate for these drawbacks, using channels or hollow fibers that can be refillable in polymers or resins for up to 30 cycles [109]. A glass fiber-reinforced composite with dual vascular system filled with epoxy resin as the healing agent and aliphatic TETA as the hardener, both separately, showed continuous healing cycles after interlaminar delaminations were achieved, resulting in full recovery (up to 100%) in fracture toughness [110]. Figure 7.15 shows the life cycle of self-healing pristine woven composite laminate with dual channels (the red channel is filled with epoxy, and the blue channel contains TETA as the hardener) in the dual vascular network [110].

Self-Healing Green Polymers and Composites 159 Pristine composite laminate

Self-healing cycle

1 Delamination damage

3 Reaction & recovery

2 Vascular release

Figure 7.15 Life cycle of a self-healing pristine woven composite laminate with dual channels (the red channel is filled with epoxy and the blue channel is TETA as the hardener) in the dual vascular network [110].

7.2.2.1 One-, Two-, or Three-Dimensional Microvascular Systems In microvascular self-healing system, hollow glass fibers (HGFs) with diameter size of 60 μm can be loaded with the appropriate healing agent. In fact, any brittle fibers could work since they would break when bent, stressed, or damaged. Typical one-, two-, or three-dimensional vascular systems are shown in Figure 7.16 [3]. One-dimensional networks can be obtained using hollow channels or fibers filled with healing agents. For example, epoxy laminate with HGFs of 15 μm diameter containing epoxy healing agent was developed by Bleay et al. [111]. The results showed 10% healing efficiency in compression strength [111]. Pang and Bond fabricated HGFs with much larger, 60 μm diameter fibers for more effective fluid filling and incorporated both HGF layers and solid glass fiber layers into the epoxy laminate [112, 113]. Their results showed up to 97% healing in strength for four-point bending tests. Two- and three-dimensional networks, shown in Figure  7.16, require more sophisticated manufacturing to ensure connectivity for self-healing applications. As fiber-composite laminate structures are produced by weaving and stacking of multiple layers, it is comparatively easy to separate and find spaces between the layers to place the filled fibers to form 3-D vascular systems. When fractures occur, networks around the cracks

160 Advanced Green Composites

100 m

1D

2D

3D

5 mm

2 mm

Figure 7.16 Vascular self-healing is organized according to the connectivity of the vascular network. One-dimensional (1-D) networks are obtained from hollow channels/ fibers filled with healing agent(s), Two- and three-dimensional (2-D and 3-D) vascular healing networks require more sophisticated manufacturing to ensure connectivity [3].

break, releasing the healing agents into those cracks. On coming into contact with the hardener or crosslinking agent, they polymerize or cross-link forming an adhesive at the damage site. Such adhesives are shown to heal the material over multiple fracture cycles, thus, extending their life. 2-D networks can integrate favorably at the interface between layers, reducing any adverse impact on mechanical strength as they maintain an increased degree of interconnectivity [114]. Bejan et al. have modeled 2-D interconnected networks to flow the healing agent or hardener in the same direction within the channels. The flow is oriented side-to-side or line-to-line to optimize matrix coverage as well as network shape and connectivity [115–117]. Williams et al. constructed self-healing 2-D networks within PVC composite sandwich panels consisting of HGFs filled with epoxy and resin. Such an arrangement showed up to 47% recovery in compression

Self-Healing Green Polymers and Composites 161 strength after 5 h at RT [118]. Toohey et al. developed 3-D vascular selfhealing epoxy composites. 10 wt% Grubbs’ catalyst particles were incorporated within the coating and the network is filled with DCDP. After 4 cycles of 4-point-bending tests and healing for 24 h at RT, it showed over 20% healing efficiency [24, 119]. In another study, a brittle epoxy coating with Grubbs’ catalyst incorporated was deposited on to a flexible epoxy substrate containing a 3-D grid network of channels having 200 μm diameter, containing a DCPD monomer as the healing agent [120]. The microvascular system, contained inside vascular fiber, offered continuous flow of healing agent, instead of one time use hollow glass fiber showing up to 97% recovery of original compression strength [120]. Carbon nanotubes have also been used as a healing agent to increase mechanical properties as well as conductivity with low loadings [121].

7.2.3 Intrinsic Self-Healing System Intrinsic self-healing materials achieve autonomous repair though inherent reversibility of non-covalent bonding of the matrix polymer. Intrinsic self-healing can be accomplished through a variety of mechanisms such as thermally reversible reactions, hydrogen bonding, ionomeric coupling, a dispersed thermoplastic phase that is able to melt, or molecular diffusion for small size of damage [3]. While reversible bonding schemes make use of the reversible nature of certain chemical reactions that can be adapted to self-healing applications, chain entanglement approaches utilize mobility at crack faces to entangle chains that span the crack surfaces. An example of intrinsic self-healing is the self-healing epoxy containing phase-separated poly(caprolactone) [123]. However, weak strength and limited crack volume of these materials limit their application to the industry as shown in Figure 7.17. Figure 7.17 shows the range of the healing efficiency, damage volume, and healing rate for the three types of self-healing approaches. In intrinsic self-healing systems, the material is inherently able to restore its integrity. However, such systems often require an external trigger such as heat, UV, and catalyst for tangible healing to occur [3]. The most widely used reversible reactions are based on Diels–Alder (DA) and retro-Diels– Alder (RDA) reactions. In this case, the monomer containing furan or maleimide form two carbon–carbon bonds and construct the polymer through DA reaction [16]. This polymer breaks down to its original monomeric units via RDA reaction when triggered by heat and reforms the polymer upon cooling or through any other conditions that were initially used to make the polymer. Another reversible reaction used in intrinsic selfhealing system is thiol/disulfide linkages [11]. The thiol-based polymers

162 Advanced Green Composites 1.0 Healing rate / damage rate

Healing efficiency, n

1.0 0.8 0.6 0.4 0.2

0.01

(a)

0.1

1.0

10

Damage volume (mm3)

0.8 0.6 0.4 0.2

0.0

100

(b)

0.2

0.4

0.6

0.8

1.2

Healing rate / damage volume

Intrinsic

Intrinsic (potential capability)

Vascular

Vascular (potential capability)

Capsule based

Figure 7.17 The range of the healing efficiency, damage volume, and rate depending on the three types of self-healing approaches [3].

have disulfide bonds that can be reversibly cross-linked through oxidation and reduction. Under reducing condition, the disulfide bridges in the polymer break down into monomers and under oxidizing conditions, the thiols in each monomer form the disulfide bond again, cross-linking the starting materials to form the polymer. Another intrinsic self-healing system can be achieved by ionomeric copolymers, which are a class of materials having ionic segments that form clusters that can act as reversible crosslinks. For example, poly(ethylene-co-methacrylic acid) copolymers with ionic segments showed their ability to restore to their original shape after healing [124]. An alternative method for achieving intrinsic self-healing is based on molecular diffusion. This system is very dependent on time and temperature, both factors controlling diffusion, and heals via surface interaction and molecular entanglement at the damaged sites [2]. As an example, self-healing via molecular diffusion and entanglement of dangling chains has been shown to occur in a weak PU gel [125].

7.2.3.1

Test Methods to Characterize Self-Healing

The primary aim of self-healing characteristic is to restore the mechanical properties reduced due to damages occurring in resins and composites. Complete filling of damaged volume such as cracks and reforming the bonds across damage surfaces can recover original properties to a great

Self-Healing Green Polymers and Composites 163 extent, if not fully. The damage modes in polymers and polymer composites causing loss of function, particularly mechanical properties, vary depending on the external stimuli. The range or extent of damage volume that can be healed differs for each approach. Figure 7.17 shows the range of the healing efficiency, damage volume, and healing rate depending on the three types of self-healing methods [3]. Most intrinsic systems can only heal small damage volumes efficiently because intrinsic re-bonding requires close proximity of the damaged surfaces. Capsule-based systems can fill small to moderate damage volumes, given a limited capsule volume fraction and vascular systems can deliver a healing agent to both small and large damage volumes. External stimuli such as impact, stress, fatigue, fracture, puncture, and corrosion are known to cause loss of function in resins and composites. By incorporating self-healing properties into various polymeric resins depending on their final applications, many researchers have proposed a variety of methods to quantify healing efficiency. The primary fracture loading conditions for self-healing specimens have been quasi-static fracture, fatigue, and impact including Mode I crack opening, Mode III tearing, mixed-mode cutting, matrix-fiber delamination, and transverse (shear mode) cracking. These loading conditions are presented in Figure 7.18 [104].

7.2.3.2

Quasi-Static Fracture Methods

Quasi-static fracture experiments are most commonly used to characterize self-healing properties in polymers and composites as they allow for controlled and predictable crack propagation including very small damage volumes [3]. For quasi-static fracture, healing efficiency (η) reduces to simply the ratio of healed to virgin fracture properties because the fracture resistance in the damaged state is zero. There are several methods in Quasistatic fracture methods. Mode I fracture method is most widely used for self-healing materials including double-cantilever beam (DCB), tapered double-cantilever beam (TDCB), compact tension (CT), single-edge notched beam (SENB), mixed-mode cutting, three- or four-point bending, and double cleavage drilled compression (DCDC). These modes are shown in Figure 7.18 [3, 18, 104, 125, 126]. Self-healing test in Mode III tearing test has been often conducted using the trouser tear test for measuring tear strength as shown in Figure 7.18(d) [104]. Figure 7.18(a, top) shows a SEM image of microcrack morphology of epoxy specimen containing DCPD microcapsules and Grubbs’ catalyst. The epoxy resins were tested using a Mode I TDCB test [18]. Figure 7.18(b) displays images of Mode I a DCDC test method of intrinsic self-healing specimen. Stable crack growth

164 Advanced Green Composites Quasi-static fracture

Top view

Healed virgin tear line Second tear line

Side view 5 m

200 m

(a) TDCD mode I plane 3 mm 10 mm

Original tear line 3 mm

10 mm B

(b) DCDC mode I

C

5 mm

(c) Mixed-mode cutting

300 m

5 mm

(d) Mode III tearing

Figure 7.18 Self-healing test methods: (a) a SEM image of microcrack plane morphology for a Mode I tapered double-cantilever beam (TDCB) self-healing epoxy specimen containing DCPD capsules and Grubbs’ catalyst [18], (b) optical images of Mode I crack growth in a double cleavage drilled compression (DCDC) intrinsic self-healing specimen based on the Diels-Alder–retro-Diels-Alder reversible bonding. Stable crack growth (left) is healed (right) after 10 h at 85 °C and 0.35 MPa clamping pressure and after 3 h at 95 °C with no clamping pressure [126], and (c) optical images of mixed-mode cutting (left) and healing (right) of an intrinsic self-healing polymer based on molecular diffusion and entanglement of dangling chains in a weak polyurethane gel. A polymer sample is stretched over a cylinder for viewing crack damage and healed morphology [125], and (d) optical image (top) and scanning electron micrograph (bottom) of a Mode III trouser tear self-healing polydimethylsiloxane (PDMS) specimen [104].

(left) was healed (right) after 10 h at 85 °C and 0.35 MPa clamping pressure and after 3 h at 95 °C with no clamping pressure [126]. Figure 7.18(c) shows optical images of mixed-mode cutting (left) and healing (right) of an intrinsic self-healing PU gel polymer. A polymer specimen is stretched over a cylinder for viewing crack damage and healed morphology [125], and Figure 7.18(d) shows optical image (top) and SEM image (bottom) of fracture surfaces after Mode III trouser tear test conducted on self-healing PDMS resin [104]. White et al. were able to demonstrate 75% recovery of virgin fracture toughness using TDCB tests of epoxy resins containing DCPD-loaded microcapsules and Grubbs’ catalyst [127]. Luo et al. evaluated specimens using three-point bending tests on SENB. Intrinsic self-healing epoxy/ poly( -caprolactone) (PCL) with 4,4 -diaminodiphenylsulfone resin resulted in up to 100% healing efficiency [123]. Cho et al. evaluated PDMSloaded microcapsules in epoxy resin using TDCB fracture test method, showing 88% healing efficiency, compared to virgin fracture toughness [33]. Wilson et al. studied epoxy vinyl ester resin consisting DCPD loaded

Self-Healing Green Polymers and Composites 165 Fatigue

500 m

50 m

5 m

Figure 7.19 SEM images of microcrack morphology for Mode I fatigue test of a TDCB self-healing epoxy specimen containing DCPD-loaded microcapsules and dispersed Grubbs’ catalyst [132].

microcapsules and dispersed Grubbs’ catalyst [128]. Their results showed to heal the epoxy vinyl ester resin with up to 30% healing efficiency using the TDCB test method [128]. Chen et al. used CT methods to characterize intrinsic self-healing furan-maleimide polymers and showed 50% healing efficiency [129]. Chen et al. also used a modified CT test method with a crack arrest hole to prevent complete failure of the specimen [130]. In their study, modified furan-maleimide polymers showed 80% recovery of fracture load. Mode I fracture toughness test was modified to a width-tapered DCB (WTDCB) for woven carbon epoxy resin composite containing DCPD microcapsules and Grubbs’ catalyst by Kessler et al. [35]. After 48 h healing period at RT, healing efficiency of up to 80% was achieved [35]. A woven glass fiber-reinforced epoxy composite containing 10 wt% epoxy-loaded UF microcapsules and dispersed latent hardener (2 wt% CuBr2(2-MeIm)4) in the composites showed 68% healing efficiency when tested using singleedge notched bending test [57]. In Mode III, the intrinsic self-healing PU gel showed restoration of about 80% of tear strength after 10 min of healing at RT after bringing together the torn crack faces [131]. Keller et al. demonstrated up to 100% recovery of tear strength in PDMS resin containing a microcapsule-based PDMS and platinum catalyst using tear strength method [104].

7.2.3.3

Fatigue Fracture Methods

Fatigue fracture is a commonly used dynamic failure mode for testing resins and composites [3]. Figure 7.19 shows SEM images of fracture surfaces of self-healing epoxy specimens containing DCPD-loaded microcapsules and Grubbs’ catalyst after Mode I fatigue TDCB test [132]. This microcapsule-based self-healing system demonstrated a fatigue life extension of

166 Advanced Green Composites 0.35 C=O Peak area ratio of (C=C)/ (C=O)

C=C 36 h 24 h 12 h 9h 6h 3h 1h 0h 0

(a)

500

1000 1500 2000 Raman shift (cm–1)

0.30 0.25 0.20 0.15 0.10

2500

(b)

0

10

20 Time (h)

30

40

Figure 7.20 In situ Raman microscopy observation: (a) the typical Raman spectra collected from the center of the fracture surface of a self-healing specimen with 15 wt% GMA-loaded microcapsules, and (b) time dependence of the characteristic Raman peak area ratio, showing the consumption of GMA during polymerization as a function of time [46].

λ = 2 for Mode I tensile fatigue test incorporating a single unloaded thermal rest period of 120 °C for 10 min. This condition resulted in 40% healing efficiency [133]. Brown et al. and Maiti et al. studied epoxy specimens containing DCPD microcapsules and Grubbs’ catalyst using Mode I fatigue TDCB [134, 135]. Their results showed 118% healing efficiency [134, 135].

7.2.3.4 Impact Fracture Methods Out-of-plane impact events can occur commonly and often result in massive damage volume from several failure modes such as puncture, delamination, and mixed-mode cracking. Healing has been quantified by restoration of compressive strength through compression-after-impact (CAI) testing [3]. Williams et al. studied 2.6  mm thick carbon fiberreinforced epoxy composites that incorporated 60 μm diameter HGF containing epoxy as the healing agent. These composites resulted in about 80% healing efficiency in compression strength [136]. Yin et al. incorporated epoxy-loaded HGF in an epoxy resin, resulting in nearly 100% efficiency using CAI tests with 1.5 J impacts [58].

7.2.3.5 Other Techniques Raman spectroscopy has been found to be a useful method that offers a live record of the surface polymerization of the released GMA on the living PMMA matrix [46]. The researchers used the C=O dependence on the peak intensity of C=C at 1638 cm−1 representing the methacrylate group as shown in Figure 7.20. The peak indicating C=C at 1638 cm−1 decreased with time. This Raman spectroscopy method served as a measure of the

Self-Healing Green Polymers and Composites 167

|Z|/W. cm2

Neat epoxy coating without scratch

1.0E+11 1.0E+10 1.0E+09 1.0E+08 1.0E+07 1.0E+06 1.0E+05 1.0E+04 1.0E+03 1.0E+02 1.0E+01 1.0E+00

Scratched tung oil capsules incorporated epoxy coating Scratched neat epoxy coating

1.0E-03 1.0E-03 1.0E-01 1.0E-00 1.0E+01 1.0E+02 1.0E+03 1.0E+04 1.0E+05

Frequency (Hz)

Figure 7.21 Electrochemical impedance spectroscopy (EIS) of self-healing coating [52].

healing polymerization kinetics, confirming that GMA flowed out of the broken microcapsules and underwent polymerization on the fracture surface of PMMA as a function of time [46]. Self-healing coating materials containing Tung oil microcapsules were characterized for their healing efficiency using electrochemical impedance spectroscopy (EIS) as shown in Figure  7.21 [52]. The corrosion resistance of scratched neat specimens and microcapsule incorporated specimens as well as unscratched virgin epoxy coated C-steel specimens were monitored over a period of two weeks. Data were obtained after three time periods (1, 7, and 14 days) after immersion by EIS and the results were compared. EIS measurements were carried out at the Ecorr potential of specimens to observe healing efficiency in the frequency range from 0.01 to 100,000 Hz [52]. Results indicated that scratched specimens compared well with the unscratched specimens indicating satisfactory self-healing.

7.3

Self-Healing Polymers from Green Sources

In the past 3–4 decades, most of the research has been developing selfhealing resins and composites as well as shells and healing agents that are petroleum based and fully synthetic resins. Examples include epoxy [5, 137], glycidyl methacrylate [45], polydimethylsiloxane [138], dicyclopentadiene [74, 139], isocyanates [1], and 4,4 -bis-methylene cyclohexane diisocyanate [140], in shell materials using urea-formaldehyde [141], melamine-formaldehyde [30], poly(dopamine) [142], melamine-ureaformaldehyde [32], or polyurea [140]. Until the last decade, there has been almost no research conducted on self-healing of green resins using green healing agents. However, this is changing fast as more and more green, and

168 Advanced Green Composites sustainable biobased resins, chemicals, and composites are being developed. The research has also been extended successfully in biomaterials used in biological applications.

7.3.1

Self-Healing Polymers in Biomaterials

Many new types of self-healing hydrogels from green sources have been successfully developed in the past couple of decades [147–159]. A common feature of dynamic polymer networks is their self-healing nature, due to their ability to undergo reversible reactions including breaking and reforming the bonds [148]. Consequently, self-healing gels can recover into their original shape after damage and further restore their original viscoelastic and tensile properties [143]. Figure 7.22 shows various approaches to synthesizing self-healing gels using physical and chemical reactions. As shown in Figure 7.22, self-healing property in hydrogels could be achieved through various physical interactions, including hydrophobic interactions, host–guest interactions, hydrogen bonds, crystallization, polymer– nanocomposite interactions, multiple intermolecular interactions, and others [144]. Chemical gels (shown at the bottom in Figure 7.22) achieve

Host-guest interaction

Hydrogen bond

Polymernanocomposite Crystallization interaction Multiple intermolecular interaction

Hydrophobic interaction

Self-healing gels

O

R1

R1

O B O OH

S R1

O

S

Phenylboronate ester complexation Disulfide bond

O

R2

R2 N

N N H

N R2

S R1

S

S

R1

O

Diels-Alder reaction

R1

Imine bond R 1 Acylhydrazone Reversible radical bond

reaction

Figure 7.22 Various approaches used to synthesize self-healing gels using physical and chemical reactions [144].

Self-Healing Green Polymers and Composites 169 their self-healing behavior through various chemical bonds, including phenylboronate complexation, disulfide bonds, imine bonds, acylhydrazone bonds, reversible radical reaction, and Diels–Alder reaction [144]. For example, a glycol chitosan-based hydrogel was developed by Yang et al. using glycol chitosan and difunctional poly(ethylene glycol) solutions. This hydrogel could recover its original shape into a complete homogeneous hydrogel with up to 97% viability after 24 h at RT [145]. A magnetic chitosan–Fe3O4 ferrofluid has been prepared by mixing with Fe3O4 nanoparticles into chitosan solution. This chitosan–Fe3O4 ferrofluid was afterwards put into with telechelic difunctional poly(ethylene glycol) (DF-PEG) solution to generate the target magnetic hydrogel. The specimen with a hole at the center recovered the original shape and viscosity after 60 min, as shown in Figure 7.23 [146]. The pH-sensitive self-healing hydrogel through complexation of a branched catechol derivatized poly(ethylene glycol) with 1,3-benzenediboronic acid was developed [151]. The pH-responsiveness of the hydrogel has been characterized, as shown in Figure 7.24, demonstrating covalent gel reactions at alkaline pH 9 and dissociation into a viscous liquid at neutral and acidic pH 3 [147]. Additionally, self-healing hydrogels are being used in steadily increasing number in biomedical applications, such as oxygen-permeable contact lenses [148], cell carriers [149], wound dressings [150], biodegradable drug delivery systems [147], and scaffolds for tissue engineering [151, 152]. For example, pH-responsive hydrogels consisting of a silk-like middle block (C2S48HC2 polypeptides) has been developed by Golinska et al. [153]. At low pH, the protein exists as monomers, but above pH 6 it readily self-assembles into long fibers. At pH 7, C2S48HC2 polypeptide chains showed selfhealing properties by recovering their original shape [153]. Self-healing chitosan-based films have shown excellent recovery from microcracks after 2 days of immersion in 0.05 M NaCl [154]. Self-healing elastomers using epoxidized natural rubber vulcanized in dicumyl peroxide have been (a) Initial

T=0

20 min

40 min

1000

Figure 7.23 Chitosan–Fe3O4 self-healing magnetic hydrogel [146].

170 Advanced Green Composites O B– O OH

HO

pH 9.0

pH 9.0

O B–

O

HO HO B

OH + HO B OH HO

pH 3.0

pH 3.0

Figure 7.24 pH-responsiveness of self-healing hydrogel from catechol derivatized poly(ethylene glycol), covalent gel reactions at alkaline pH 9 and dissociation into a viscous liquid at neutral and acidic pH 3 [147].

studied for the ballistic vest applications. The damage initiated by T-peel test was observed to heal the hole after 72 h healing period at 60 °C [155]. Poly(lactide-co-glycolide) microcapsules, containing ricin toxoid or fluorescein isothiocyanate-labeled bovine serum albumin were prepared by a water-in-oil-in-water emulsion solvent extraction procedure with a high encapsulation efficiency of up to 94% [90]. Polysaccharide-based selfhealing N-carboxyethyl chitosan had 94% healing efficiency in fracture toughness test after 24 h healing period at RT [156].

7.3.2 Self-Healing Green Resins and Green Composites Just in the past couple of years, the self-healing technology has been extended to green polymers. This includes microcapsule-based system developed for self-healing green polymers such as plant-based proteins and starches using green healing agent, green solvent, and green shell material. For example, self-healing soy protein isolate (SPI)-based “green” thermoset

Self-Healing Green Polymers and Composites 171

SPI

NH2

c c c

GA

Crosslinking N

CH

H2N

(a)

(b)

Soy protein encapsulated with PLGA (SPI- PLGA) Glutaraldehyde

N

CH

CH2 CH2

NH2

CH2

CH2

CH2

CHO

CH2

CH

NH2

N

CHO

H2N N

CH

CH2

CH2

C

CHO CH

CH2

CH2

C

CH n

CH2

CH2 CH2

CHO

(c)

Figure 7.25 Self-healing reaction and bridging zone of soy protein isolate-loaded poly(d,l-lactide-co-glycolide) (SPI-PLGA) microcapsules in SPI resin: (a) the cross-linking reaction between SPI and glutaraldehyde, (b) fracture surface of SPI resin containing SPIPLGA microcapsules, and (c) bridging zone healed from released SPI from microcapsules between two fracture surfaces in SPI resin [157].

resin was developed by Kim and Netravali using PLGA microcapsules containing SPI, as a crack healant. The self-healing efficiency was investigated using Mode I fracture toughness test for resins containing 15 wt% microcapsules and 12 wt% glutaraldehyde as the cross-linking agent. The healing resulted by bridging the two fracture surfaces and showed selfhealing efficiency of up to 48%. The schematic, cross-linking reaction, and crack bridging (SEM) are shown in Figure 7.25 [157]. They expanded their research to evaluate the effect of various parameters such as PLGA or PVA concentration and homogenization speed on microcapsule size and distribution as well as self-healing properties of SPI resins. Their results showed increased healing efficiency of up to 53% using higher PVA concentration of 5% [158]. They explained that higher PVA concentration increases the bonding between the microcapsules and the resin ensuring the fracture of microcapsules in the microcrack region. Wertz et al. showed that PLA films

172 Advanced Green Composites

Entrance

Exit

Figure 7.26 Optical microscopy images of bullet entrance and exit zone after self-healing of ethylene/methacrylic acid copolymer resin [161].

containing 2.5 wt% catalyst and 7.5 wt% DCPD loaded microcapsules, of around 31.5 μm diameter, increased fatigue toughness by up to 84% compared to virgin PLA films [139]. Polysaccharide-based self-healing hydrogels developed using the reaction of N-carboxyethyl chitosan and adipic acid dihydrazide with oxidized sodium alginate showed 94% healing efficiency as measured by their rheological recovery [156]. Xiang et al. developed self-healing vulcanized rubber using CuCl2 as the catalyst, which was capable of re-arranging disulfide linkages within the microcrack in the vulcanized rubber, recovering the original tensile properties after compression molding at 110 °C [159]. Fully bio-based self-healing epoxidized soybean oil resin was developed by Altuna et al. [160]. They used an aqueous citric acid solution as the cross-linking agent. Soybean oil-based resin recovered its original fracture stress and strain properties, as assessed by stress relaxation and lap-shear tests, after compression molding at 160 °C. According to them, this finding was a result of the residual epoxy groups from soybean oil reacting with carboxylic groups from the citric acid and providing additional cross-linking within the resin [160]. The self-healing behavior of ethylene/methacrylic acid copolymer resin with 15–50 wt% loading of epoxidized natural rubber was investigated. The self-healing behavior at entrance and exit regions of a bullet was observed after ballistic puncture tests as shown in Figure 7.26 [161]. The self-healing behavior of blends of cis-1,4-polyisoprene and epoxidized natural rubber resin with dicumyl peroxide (DCP) as a cross-linking

Self-Healing Green Polymers and Composites 173

Apparent cross-link density (1/Q)

0.24 0.22 0.2 0.18 0.16 0.14 0.12 Self-healing No healing

0.1 0.2

0.5

0.8 1 DCP content (wt%)

2

3

Figure 7.27 The self-healing behavior of blends of cis-1,4-polyisoprene and epoxidized natural rubber with dicumyl peroxide as a cross-linking agent [162].

agent for rubbers was investigated by Burfield et al. using T-peel tests [162]. A higher epoxidation level was found to enhance self-healing by way of cross-linking reactions between the epoxide groups and peroxide, at RT, as shown in Figure 7.27 [162]. Kim and Netravali also developed self-healing waxy maize starch-based thermoset resin [163]. In this study, self-healing was achieved using waxy maize starch-loaded PLGA microcapsules. Starch resin containing 20 wt% microcapsules showed self-healing efficiency of up to 51% in fracture stress and 66% in fracture toughness after 24 h of healing at RT. Kim and Netravali also extended the self-healing technology to fully green composites made using SPI-based resin and microfibrillated cellulose (MFC, 10% by wt) [164]. The PLGA microcapsules contained SPI as the healant. At 15 wt% microcapsules, the composites showed about 27% healing after 24 h of healing at RT. The reason for the lower self-healing efficiency, compared to pure resin, was explained as because the healing occurs only in the resin; whereas, the broken MFC fibrils that reinforce the resin cannot be healed.

7.4 Summary and Prospects An overview of various self-healing concepts developed for polymers for use as resins and coatings as well as polymeric composites published over

174 Advanced Green Composites the past 20 years is presented in this chapter. This is an emerging and fascinating field of research, and successful self-healing technologies could significantly extend the working life and safety of the polymers and their composites for a broad range of applications. To date, most of the development has been in conventional, petroleum-based polymers, and composites. In addition, the self-healing polymeric materials have been largely based on mimicking biological healing. Despite the significant advancements made, there is still a long way to go before even the simplest biological healing mechanism can be replicated using these synthetic materials for real life. One main reason is that the biological healing occurs in live tissues that can be replenished continuously whereas in the case of polymeric materials, this is not true. Never-the-less, past 2–3 decades have seen great progress made in this area and yet, the use of self-healing materials is limited even for the synthetic polymers. These materials are not biocompatible and hence cannot be easily used as bone, skin, and joint replacements. However, it is certain that future developments in self-healing resins and composites will produce a new generation of materials for structural applications. It can be anticipated that the self-healing technology will someday evolve beyond the current state to procedures that replicate biomimetic healing abilities by combining all related fields including biomaterials, civil engineering, food engineering, drugs delivery systems, textiles/fibers, and cosmetics. Application of self-healing coatings could also improve corrosion protection, thereby extending the useful life of metallic structures. A wide ranging engineering structures from cars to aircrafts and from factory equipment to household appliances can be effectively protected from rust by the self-healing coating systems while at the same time retaining their aesthetics. However, further investigation of these coating to assess their benefits is recommended. The present review covers only the micro/nanocapsule-based system that responds to mechanical degradation or microcrack propagation. Much more research in terms of new material synthesis and application of new micro/nanocapsule family that are sensitive to pH, temperature, environmental conditions changes, and their dispersion to coatings may be considered. Research in sustainable or renewable polymers derived from plants is an emerging field that is expected to grow exponentially as the supplies of petroleum dwindle and, at the same time, concerns over the land, air, and water pollution created by the polymer waste grow. The situation has already reached a critical point. As new green polymers and composites are developed, the self-healing technologies developed for conventional polymers and composites can be easily transferred to them as well, extending their life span and making them more acceptable in mainstream

Self-Healing Green Polymers and Composites 175 applications. The self-healing coatings when applied to green polymers and composites could help propel their use as well.

Acknowledgements The authors would like to acknowledge the partial financial support of National Institute of Food and Agriculture (NIFA), U.S. Department of Agriculture (Multistate Research Project S-1054 under 1004862).

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182 Advanced Green Composites 104. M. W. Keller, S. R. White and N. R. Sottos, A Self-Healing Poly (Dimethyl Siloxane) Elastomer. Adv. Function. Mater. 17, 2399–2404, 2007. 105. S. H. Cho, S. R. White and P. V. Braun, Self-healing polymer coatings. Adv. Mater. 21, 645–649, 2009. 106. D. S. Xiao, Y. C. Yuan, M. Z. Rong and M. Q. Zhang, Self-healing epoxy based on cationic chain polymerization. Polymer. 50, 2967–2975, 2009. 107. Q. Li, A. K. Mishra, N. H. Kim, T. Kuila, K.-t. Lau and J. H. Lee, Effects of processing conditions of poly (methylmethacrylate) encapsulated liquid curing agent on the properties of self-healing composites. Compos. Part B. 49, 6–15, 2013. 108. A. Yabuki, H. Yamagami and K. Noishiki, Barrier and self-healing abilities of corrosion protective polymer coatings and metal powders for aluminum alloys. Mater. Corros. 58, 497–501, 2007. 109. N. K. Guimard, K. K. Oehlenschlaeger, J. Zhou, S. Hilf, F. G. Schmidt and C. Barner-Kowollik, Current Trends in the Field of Self-Healing Materials. Macromol. Chem. Phys. 213, 131–143, 2012. 110. J. F. Patrick, K. R. Hart, B. P. Krull, C. E. Diesendruck, J. S. Moore, S. R. White and N. R. Sottos, Continuous Self-Healing Life Cycle in Vascularized Structural Composites. Adv. Mater. 26, 4302–4308, 2014. 111. S. Bleay, C. Loader, V. Hawyes, L. Humberstone and P. Curtis, A smart repair system for polymer matrix composites. Compos. Part A. 32, 1767–1776, 2001. 112. J. W. C. Pang, and I. P. Bond, ‘Bleeding composites’: damage detection and self-repair using a biomimetic approach. Compos. A. 36, 183–188, 2005. 113. J. W. C. Pang, and I. P. Bond, A hollow fiber reinforced polymer composite encompassing self-healing and enhanced damage visibility. Compos. Sci. Technol. 65, 1791–1799, 2005. 114. H. R. Williams, R. S. Trask, P. M. Weaver and I. P. Bond, Minimum mass vascular networks in multifunctional materials. J. R. Soc. Interface. 5, 55–65, 2008. 115. H. Zhang, S. Lorente and A. Bejan, Vascularization with trees that alternate with upside-down trees. J. Appl. Phys. 101, 094904, 2007. 116. S. Lorente, and A. Bejan, Vascularized smart materials: designed porous media for self-healing and self-cooling. J. Porous Media. 12, 1–18, 2009. 117. K. M. Wang, S. Lorente, and A. Bejan, Vascularized networks with two optimized channel sizes. J. Phys. D. 39, 3086–3096, 2006. 118. H. R. Williams, R. S. Trask and I. P. Bond, Self-healing composite sandwich structures. Smart Mater. Struct. 16, 1198–1207, 2007. 119. K. S. Toohey, N. R. Sattos, J. A. Lewis, J. S. Moore and S.R. White, Self-healing materials with microvascular networks. Nat. Mater. 6, 581–585, 2007. 120. J. Pang and I. Bond, ‘Bleeding composites’—damage detection and self-repair using a biomimetic approach. Compos. Part A. 36, 183–188, 2005. 121. J. A. Syrett, C. R. Becer and D. M. Haddleton, Self-healing and self-mendable polymers. Polym. Chem. 1, 978–987, 2010.

Self-Healing Green Polymers and Composites 183 122. G. Williams, R. Trask and I. P. Bond, A self-healing carbon fiber reinforced polymer for aerospace applications. Compos. A. 38, 1525–1532, 2007. 123. X. Luo, R. Ou, D. E. Eberly, A. Singhal, W. Viratyaporn and P. T. Mather, A thermoplastic/thermoset blend exhibiting thermal mending and reversible adhesion. ACS Appl. Mater. Interfaces. 1, 612–620, 2009. 124. S. J. Kalista, T. C. Ward and Z. Oyetunji, Self-healing of poly(ethylenecomethacrylic acid) copolymers following projectile puncture. Mech. Adv. Mater. Struct. 14, 391–397, 2007. 125. M. Yamaguchi, S. Ono and M. Terano, Self-repairing property of polymer network with dangling chains. Mater. Lett. 61, 1396–1399, 2007. 126. T. A. Plaisted and S. Nemat-Nasser, Quantitative evaluation of fracture, healing and re-healing of a reversibly cross-linked polymer. Acta Materialia. 55, 5684–5696, 2007. 127. S. R. White, N. R. Sottos, P. H. Geubelle, J. S. Moore, and M. R. Kessler, Autonomic healing of polymer composites. Nature. 409, 794–797, 2001. 128. G. O. Wilson, J. S. Moore, S. R. White, N. R. Sottos and H. M. Andersson, Autonomic healing of epoxy vinyl esters via ring opening metathesis polymerization. Adv. Funct. Mater. 18, 44–52, 2008. 129. X. Chen, M. A. Dam, K. Ono, A, Mal, H. Shen and S. R. Nutt, A thermally remendable cross-linked polymeric material. Science. 295, 1698–1702, 2002. 130. X. Chen, F. Wudl, A. K. Mal, H. B. Shen and S. R. Nutt, New thermally remendable highly cross-linked polymeric materials. Macromolecules. 36, 1802–1807, 2003. 131. M. Yamaguchi, S. Ono and K. Okamoto, Interdiffusion of dangling chains in weak gel and its application to self-repairing material. Mater. Sci. Eng. B. 162, 189–194, 2009. 132. E. N. Brown, S. R. White, and N. R. Sottos, Fatigue crack propagation in microcapsule-toughened epoxy. J. Mater. Sci. 41, 6266–6273, 2006. 133. M. Zako and N. Takano, Intelligent material systems using epoxy particles to repair microcracks and delamination damage in GFRP. J. Intell. Mater. Syst. Struct. 10, 836–841, 1999. 134. E. N. Brown, S. R. White, and N. R. Sottos, Retardation and repair of fatigue cracks in a microcapsule toughened epoxy composite—Part II: In situ selfhealing. Compos. Sci. Technol. 65, 2474–2480, 2005. 135. E. N. Brown, S. R. White, and N. R. Sottos, Retardation and repair of fatigue cracks in a microcapsule toughened epoxy composite. 1. Manual infiltration. Compos. Sci. Technol. 65, 2466, 2005. 136. G. J. Williams, I. P. Bond and R. S. Trask, Compression after impact assessment of self-healing CFRP. Compos. A. 40, 1399, 2009. 137. T. S. Coope, U. F. J. Mayer, D. F. Wass, R. S. Trask and I. P.Bond, Self-Healing of an Epoxy Resin Using Scandium(III) Triflate as a Catalytic Curing Agent. Adv. Function. Mater. 21, 4624–4631, 2011.

184 Advanced Green Composites 138. S. H. Cho, H. M. Andersson, S. R. White, N. R. Sottos and P. V. Braun, Polydimethylsiloxane-Based Self-Healing Materials. Adv. Mater. 18, 997– 1000, 2006. 139. J. T. Wertz, T. C. Mauldin and D. J. Boday, Polylactic Acid with Improved Heat Deflection Temperatures and Self-Healing Properties for Durable Goods Applications. ACS App. Mater. Inter. 6, 18511–18516, 2014. 140. D. Sun, H. Zhang, X.-Z. Tang and J. Yang, Water resistant reactive microcapsules for self-healing coatings in harsh environments. Polymer. 91, 33–40, 2016. 141. B. Blaiszik, N. Sottos and S. White, Nanocapsules for self-healing materials. Compos. Sci. Technol. 68, 978–986, 2008. 142. A. Jones, C. Watkins, S. White and N. Sottos, Self-healing thermoplastictoughened epoxy. Polymer. 74, 254–261, 2015. 143. U. Gulyuz and O. Okay, O., Self-healing polyacrylic acid hydrogels. Soft Matter. 9, 10287–10293, 2013. 144. Z. Wei, J. H. Yang, J. Zhou, F. Xu, M. Zrínyi, P. H. Dussault, Y. Osada and Y. M. Chen, Self-healing gels based on constitutional dynamic chemistry and their potential applications. Chem. Soc. Rev. 43, 8114–8131, 2014. 145. B. Yang, Y. Zhang, X. Zhang, L. Tao, S. Li and Y. Wei, Facilely prepared inexpensive and biocompatible self-healing hydrogel: a new injectable cell therapy carrier. Polym. Chem. 3, 3235–3238, 2012. 146. Y. Zhang, B. Yang, X. Zhang, L. Xu, L. Tao, S. Li, and Y. Wei, A magnetic selfhealing hydrogel. Chem. Commun. 48, 9305–9307, 2012. 147. P. Gupta, K. Vermani and S. Garg, Hydrogels: from controlled release to pHresponsive drug delivery. Drug Discovery Today. 7, 569–579, 2002. 148. J. W. Kwon, Y. K. Han, W. J. Lee, C. S. Cho, S. J. Paik, D. I. Cho, J. H. Lee and W. R.Wee, Biocompatibility of poloxamer hydrogel as an injectable intraocular lens: a pilot study. J. Cataract Refractive Surgery. 31, 607–613, 2005. 149. L. N. Novikova, A. Mosahebi, M. Wiberg, G. Terenghi and J. O. Kellerth, Alginate hydrogel and matrigel as potential cell carriers for neurotransplantation. J. Biomed. Mater. Res. Part A. 77, 242–252, 2006. 150. B. Balakrishnan, M. Mohanty, P. Umashankar and A. Jayakrishnan, Evaluation of an in situ forming hydrogel wound dressing based on oxidized alginate and gelatin. Biomaterials. 26, 6335–6342, 2005. 151. J. L. Drury and D. J. Mooney, Hydrogels for tissue engineering: scaffold design variables and applications. Biomaterials. 24, 4337–4351, 2003. 152. T. R. Hoare and D. S. Kohane, Hydrogels in drug delivery: progress and challenges. Polymer. 49, 1993–2007, 2008. 153. M. D. Golinska, M. K. Włodarczyk-Biegun, M. W. Werten, M. A. C. Stuart, F. A. de Wolf and R. de Vries, Dilute self-healing hydrogels of silk-collagen-like block copolypeptides at neutral pH. Biomacromolecules. 15, 699–706, 2014. 154. M. Zheludkevich, J. Tedim, C. Freire, S. Fernandes, S. Kallip, A. Lisenkov, A. Gandini, and M. G. S. Ferreira, Self-healing protective coatings with

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8 Transparent Green Composites Antonio Norio Nakagaito1*, Yukiko Ishikura2 and Hitoshi Takagi1 1

Graduate School of Technology, Industrial and Social Sciences, Tokushima University, Tokushima, Japan 2 Research Institute for Sustainable Humanosphere, Kyoto University, Uji, Kyoto, Japan

Abstract Optically transparent composites have been developed in recent years to substitute glass, which is currently being used as substrates in electronic devices. Flexible electronics is now a common technology employed in gadgets that are ever-present in our daily lives. Among them, rigid electronic displays are about to become flexible in the coming years. Plastics are expected to be the material of choice for substrates due to their intrinsic flexibility and optical qualities, with the only inconvenience of being prone to thermal expansion. The compatibility between the thermal expansions of the substrate and the deposited active layers is critical to avoid damages when subjected to thermal cycles during device manufacture and use. Thermal expansion of plastics can be restrained by reinforcing them with nanofibers without appreciable reduction in their transparency. Good gas barrier properties of these transparent films have also brought the possibility to use them in packaging. Nanofibers are available in large amounts in nature as cellulose and chitin that are produced by plants and animals. In this chapter, a variety of transparent composites, some fully green and others partially green, are presented and discussed. Keywords: Transparency, composite, green, cellulose, chitin, nanofiber, thermal expansion, gas barrier

8.1 Introduction Flexible displays are made possible by the new organic light-emitting diode (OLED) panels that are slowly replacing the older but time proven *Corresponding author: [email protected] Anil Netravali (ed.) Advanced Green Composites, (187–210) © 2018 Scrivener Publishing LLC

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188 Advanced Green Composites liquid crystal displays (LCD). The use of flexible plastic substrates instead of rigid glass can produce tough, thin, and lightweight panels that flex, bend and conform to desired shapes as wearable accessories or can be simply rolled for easy portability. Continuous roll-to-roll processing could also be adopted in mass production of display units. However, for this to be successful, plastic substrates must offer properties such as high transparency, surface smoothness, solvent resistance, barrier capabilities, thermal and dimensional stabilities, and reduced thermal expansion, similar to those offered by glass. Bottom-emissive displays require a total light transmittance higher than 85% over the visible spectrum (400 to 800 nm) with a haze of less than 0.7%, since images are viewed through the substrate [1]. The barrier capability and solvent resistance can be achieved by applying a hard coating, while the surface smoothness is achieved by a scratch-resistant coating. As the thermal and dimensional stabilities are properties inherent to the material, the right type of plastic must be chosen to resist the high temperatures involved in the deposition of coatings and active layers needed in display manufacturing. However, the coefficient of thermal expansion (CTE) below 20 10 6 K 1 is required to avoid any mismatch in thermal expansion between the deposited device layers and the substrate that would cause any damage during a thermal cycle [1]. From the atomistic view of materials, thermal expansion and Young’s modulus depend on the depth of the atomic bond energy function. The relationship between the two properties, though, is inversely proportional, making flexible plastics having low modulus exhibit intrinsically high CTEs. One sought out way to reduce the CTE of plastics while preserving its flexibility, is reinforcing them with fibers that possess low CTE. However, as micro-sized fibrous elements reduce the transparency due to light scattering, unless the refractive indexes of fibers and plastic matrix are exactly matched, there is a need to use nanofibers with lateral dimensions much smaller than the wavelength of the visible light. As the visible spectrum resides in the range of 400 to 800 nm, the reinforcing elements must have dimensions below 100 nm. Such nanofibers are already available abundantly in nature as fully renewable or ‘green’ resources, mostly based on plants and other living creatures. Typical examples of such reinforcing materials are cellulose and chitin found mainly as the framework of higher plants and the exoskeleton of arthropods, respectively. This chapter aims to present an overview of the history and recent developments related to green transparent composites.

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8.2 Cellulose Nanofiber-Based Composites and Papers 8.2.1 Bacterial Cellulose-Based Composites The first example of an optically transparent composite reinforced with cellulose nanofibers (CNF) was reported by Yano et al. in 2005. The nanofibers consisted of cellulose of animal origin, known as bacterial cellulose (BC) [2]. These nanofibers are synthesized and spun by certain species of bacteria consisting of pellicles formed by a continuous network of dimensionally uniform CNFs less than 100 nm in diameter. As BC pellicles are obtained originally as water containing hydrogels, they were compressed to squeeze out the excess water and dried to make thin paper-like sheets. These sheets were then combined with transparent acrylic, epoxy, and phenolic resins by impregnation to produce transparent composites. Even with high nanofiber contents of 60 to 70 wt%, the total light transmittance in the visible spectrum from 500 to 800 nm surpassed 80%. The light transmittance of these composites was reduced by less than 10 percentage points relative to the neat resins, despite the high fraction of nanofibers. In microcomposites, the light transmittance depends essentially on the refractive index difference between the reinforcing and matrix phases, as their interfaces scatter most of the incident light. The refractive index of CNFs is 1.618 lengthwise and 1.544 in the transverse direction at a wavelength of 587.6 nm and 23 °C. One of the resins used in the study, phenol formaldehyde, had a refractive index of 1.483 under the same measurement conditions. Despite the significant difference in refractive indices, light scattering was considerably attenuated due to the much smaller size of the nanofibers (~60 nm) relative to the wavelength of the incoming light, greater than 400 nm. The effect of the refractive index mismatch between resin and CNFs was shown to be minimal in a later study by Nogi et al. [3]. The total and regular transmittances of BC-reinforced acrylic resins were higher than 85% and 75% respectively, at a wavelength of 590 nm when the refractive index of the resin varied from 1.492 to 1.636. The transmittance was also highly stable with respect to variations in temperature which causes changes in the refractive index of the resin, from room temperature to 80 °C, as shown in Figure 8.1. The reduction in optical transmittance was linearly proportional to the nanofiber content, however, it dropped by only 13.7 percentage points despite the high 66 wt% nanofiber load [4]. The CTE, however, was reduced significantly from 86 × 10 6 K 1 for neat acrylic resin to 15 × 10 6 K 1 for a nanofiber fraction of 30 wt%, and to an even lower value of 10 × 10 6 K 1 for 50 wt% nanofiber content. In addition, these

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Figure 8.1 Stability of the optical transmittance of BC nanofiber composites compared to the variations in refractive index of matrix resins (three types shown) due to temperature change. [3]

fibrous nanocomposites were many times stronger than engineered plastics, reaching a tensile strength up to 325 MPa and Young’s modulus of 20–21 GPa. The CTE of BC-reinforced phenol formaldehyde was as low as that of silicon crystal at 3 × 10 6 K 1 [2]. An even lower CTE was achieved by reinforcing an already low stiffness (Young’s modulus) transparent resin with just 5 wt% BC. The resin had its

Transparent Green Composites 191 modulus increased from 25 MPa to 355 MPa when made into composites using BC, still low enough to be flexible and foldable. The in-plane CTE, however, was reduced to a very low value of 4 × 10 6 K 1 [5]. In general, from the atomistic viewpoint modulus and CTE are inversely proportional, but the anisotropic structure of BC makes these composites behave differently. Pellicles of BC are made of multiple layers of nanofiber networks weakly connected through the thickness but stiffer in-plane layers networked by branched, randomly oriented, and continuously connected nanofibers. Due to this unique morphology, even low BC content was able to restrain the small stresses produced by the thermal expansion of the low modulus matrix, whereas the bulk of the expansion happened in the thickness direction. These characteristics made BC-acrylic resin composites highly flexible and foldable with reduced in-plane CTE. As cellulose is highly hygroscopic as most biomaterials, Nogi et al. and Ifuku et al. employed acetylation treatment to replace the exposed hydroxyl groups on BC nanofibers surface by acetyl groups [6, 7]. A BC composite with about 60 wt% nanofiber content exposed to 55% relative humidity at room temperature had 3.12% moisture content whereas acetylated BC-acrylic resin composite had the moisture absorption reduced to 1.33%, more than 50% reduction. The CTE of BC sheets was reduced from 3 × 10 6 K 1 of untreated BC to 0.8 × 10 6 K 1 after acetylation, but the Young’s modulus was also reduced from 23.1 to 17.3 GPa. The regular light transmittance of BC-acrylic resin composite was increased at a wavelength around 400 nm with acetylation, while for the rest of the spectrum it was unchanged. Scanning electron microscopy (SEM) images showed that acetylated BC nanofibers remained as individualized fibers in contrast to more compacted untreated BC, suggesting that acetyl groups on the surface of nanofibers prevented their aggregation, minimizing light scattering at shorter wavelengths. Acetylation contributed positively to the degradation by heat while the loss of transmittance was less notable, when composites were exposed to 200 °C for a few hours.

8.2.2 CNF-Based Composites Since BC nanofibers are synthesized and directly secreted by certain species of bacteria, they are dimensionally uniform and highly crystalline, being the ultimate CNF morphology known. However, bacteria cultivation processes are time consuming and the carbon and nitrogen sources can be expensive. As a result, BC production costs are generally high. Another source for obtaining CNFs are plant fibers, the only caveat being the need for them to be properly extracted. Differently from BC which is produced

192 Advanced Green Composites in a bottom-up way, these nanofibers form the framework of complex biocomposites comprising the cell wall of plant fibers. The matrix or bonding substances present in these fibers such as lignin and hemicelluloses have to be chemically removed and the nanofibers extracted in individualized form to be employed in composites. Cellulose structures extracted from wood pulp fibers by refining and high-pressure homogenization, known as microfibrillated cellulose (MFC), were additionally fibrillated by a mechanical grinding treatment to obtain more dimensionally uniform nanofibers in the range of 50 to 100 nm [8–10]. The original MFC is partially nanofibrillated, containing a wide distribution of fibril diameters up to a few micrometers and reduces the transparency of final composites. The mechanically ground nanofibers were dispersed in water, vacuum filtered and dried to produce thin paper-like sheets. Sheets were then impregnated with acrylic resin and UV light-cured. The total light transmittance of these composites at 70 wt% of nanofibers was 70% at 600 nm. The transmittance was 20 percentage points lower than that of neat acrylic resin. The reduction in fibril diameter by grinding was significantly effective as the original MFC-reinforced acrylic resin at 62 wt% MFC content exhibited a transmittance of only 30%. The CTE of the acrylic resin was reduced from 86 10 6 K 1 to just 17 10 6 K 1, while the Young’s modulus was increased to 7 GPa. This was lower than that obtained by BC reinforcement but it made the composite more flexible. A finer CNF morphology was then developed by Abe et al. by a new extraction protocol that produced more dimensionally uniform nanofibers about 15 nm in diameter by a single pass through the grinder [11]. The less intensive but effective and efficient fibrillation avoided the decrease in reinforcing and thermal expansion restriction abilities due to decrease in crystallinity and degree of polymerization of cellulose. The key point was to keep the fibers wet during the whole process of matrix substances (lignin, hemicelluloses, etc.) removal and fibrillation, thus, eliminating the formation of hydrogen bonds between nanofibers that would restrict their proper individualization. A proof of concept of flexible displays based on such plant derived CNF-reinforced transparent plastic substrates was reported by Okahisa et al. [12]. The nanofibers extracted by Abe’s method were blended with a low Young’s modulus resin to form composites and an OLED layer was deposited on them. The flexible display obtained using this composite is shown in Figure 8.2. A chemical method to extract an even finer CNF morphology was developed by Saito et al. by oxidizing the surface of nanofibers by a 2,2,6,6-tetramethylpiperidine-1-oxyl (TEMPO) radical catalyzed process [13]. In this process, part of the hydroxyl groups on the CNFs’ surface are

Transparent Green Composites 193

Figure 8.2 Organic light emitting diode deposited on a cellulose nanofiber-reinforced composite film. [12]

converted into negatively charged sodium carboxylate groups that loosen their mutual adhesion by electrical repulsion. A mild agitation by an ordinary blender then individualized the nanofibers of 3 to 5 nm in diameter. Fujisawa et al. demonstrated the effectiveness of TEMPO-mediated oxidized cellulose nanofibers (TOCN) in reinforcing polystyrene (PS), especially with nanofibers with higher aspect ratios [14]. The addition of 10 wt% TOCN with aspect ratio in the range of 310, resulted in the reduction of CTE of PS from 116 10 6 K 1 to 63.7 10 6 K 1 with light transmittance of 86%, compared to 89% obtained for neat PS at 600 nm. At the same time Young’s modulus and tensile strength were increased by 91% and 18% relative to the neat resin, respectively. To explore the possibility of making transparent composites from micro-sized cellulose fibers, Yano et al. impregnated acrylic resin into an acetylated sheet of paper comprised of pulp fibers [15]. At the fiber content of 28 wt% the light transmittance achieved was 67.6% at 600 nm wavelength. The CTE of the acetylated pulp composite was reduced to 8.3 10 6 K 1 from 213 10 6 K 1 obtained for neat resin. However, the untreated pulp composite resulted in even lower CTE of 3.6 10 6 K 1. This may be due to the enhanced number of interfibrillar hydrogen bonds making the whole cell wall of untreated fibers stiffer than acetylated pulp, as the latter have less hydroxyl groups to mutually bond the nanofibers. Although not completely fruitful in terms of the transparency achieved, this study clearly hinted at the possibility of producing transparent composites without the need of the costly nanofibrillation treatment. In a greener approach to make composites, Wu et al. produced TOCN/ montmorillonite nanoplatelet-layered composites that had Young’s modulus, tensile strength and strain at fracture increased by addition of just 5 wt% of montmorillonite to 18 GPa, 509 MPa, and 7.6%, from 11.6 GPa, 210 MPa, and around 3% of the pure TOCN film, respectively [16]. The work

194 Advanced Green Composites of fracture or toughness was increased six-fold to 25.6 MJ m 3 from that of the TOCN film, much higher than those obtained with nanocellulose films. The density was slightly higher at 1.99 g cm 3 and light transmittance was around 80%. The oxygen permeability of 1 wt% montmorrilonite-added TOCN film at 0% RH was below 0.007 mL greek m 2 day 1 kPa 1. Fujisawa et al. grafted TOCN surfaces with polyethylene glycol chains via ion-exchange treatment, to obtain nano-dispersed suspension in chloroform [17]. The suspension was mixed with poly(L-lactide) acid (PLLA) previously dissolved in chloroform, to obtain films by casting. The mechanical properties of the composite films were increased significantly by merely 1 wt% addition of TOCN. The tensile strength, modulus and work of fracture were increased by 40, 26, and 52% over those of neat PLLA. Previous composites based on CNF-reinforced PLLA required nanofiber addition of over 10 wt% to achieve similar mechanical properties. The light transmittance attained at 600 nm was 89%, very close to that of neat PLLA which was 90%. Another green transparent composite reported by Cai et al. was produced by reinforcing polyvinyl alcohol with electrospun cellulose acetate nanofibers. Nanofibers had an average diameter of about 700 nm but the light transmittance of composites stayed around 75% for a nanofiber content of 37 wt% [18].

8.2.3 Transparent Nanopapers A truly green transparent material was realized not by a composite but by plain paper [19]. It was made of the same chemical substance of ordinary paper, i.e. cellulose, with the only difference being the physical size of the CNFs extracted by the process developed by Abe et al. [11]. An aqueous suspension of nanofibers was dewatered and dried as usual. As the density of the paper reached 1.53 g cm 3, a value close to the 1.59 g cm 3 of the cellulose crystal, the interstitial voids between the nanofibers were mostly eliminated [20]. Although the paper was initially translucent, instead of impregnating with a transparent resin, eliminating the surface roughness by polishing resulted in light transmittance of 71.6% at a wavelength of 600 nm. The CTE of the paper was as low as 8.5 10 6 K 1, and the mechanical properties were as high as the best nanofiber papers at 13 GPa of tensile modulus and 223 MPa of tensile strength. Another approach to reduce surface roughness was to coat the nanofiber paper with transparent resin films in a potentially continuous roll-to-roll fabrication [21]. Another example of nanopaper was reported by Nogi et al. who produced all cellulose films from 15 nm wide nanofibers extracted by a water jet system equipped with a ball-collision chamber [22]. Films produced

Transparent Green Composites 195 by casting were highly transparent with total light transmittance of 90.1% in the range of 450 to 800 nm wavelength, matching the theoretically predicted value. The haze also showed to be quite low at 4.1%. These values are quite close to those of PET films at 89% total transmittance and 4.5% haze. While PET films heated at 150 °C for 1 hour have the haze increased to 14% due to increasing surface roughness, all CNF films had both total transmittance and haze unaltered during and after heating due to the high thermal stability of cellulose. Films of TOCN have high gas barrier properties besides optical transparency. Fukuzumi et al. explained that the extremely low gas permeability of TOCN films is due to the high crystallinity of nanofibers but also due to the small pore size created in the interstitial spaces between nanofibers [23]. Positron annihilation lifetime spectroscopy determined an average pore size of about 0.47 nm, which is slightly larger than the kinetic diameter of oxygen gas. This confirmed that the high transparency of TOCN films is due not only to the 3 to 4 nm diameter of the nanofibers which are 1/100th of 400 nm wavelength, smallest in visible spectrum, but also due to the extremely small pore sizes that avert light scattering. Sun et al. compared TOCN films with films of sulfuric acid-hydrolyzed cellulose nanocrystals (SACN) [24]. Both films were highly transparent with transmittance of over 90% at 600 nm, but TOCN films had higher transparency over the whole visible spectrum. The tensile strengths were 237 MPa and 206 MPa, while the moduli were 6.6 GPa and 6 GPa for TOCN and SACN, respectively. The average CTEs were 8.41 and 8.38 10 6 K 1 for TOCN and SACN films, respectively. Fukuzumi et al. reported that films of TOCN extracted from softwood had transmittance of about 90% while those extracted from hardwood had a lower transmittance of 78%, both measured at 600 nm [25]. The probable reason was attributed to the presence of residual hemicellulose glucomannan in softwood containing C6 primary hydroxyls convertible to sodium carboxylate that facilitates nanofibrillation, whereas hardwood contains predominantly xylan with no C6 hydroxyl groups. The CTE of TOCN films is 2.7 10 6 K 1, a value lower than that of glass at 9 10 6 K 1.

8.2.4 All Cellulose Transparent Composites An additional category of transparent composites are the self-reinforced composites derived from pure cellulose. Composites consisting solely of cellulose were produced using partially carboxymethylated softwood kraft pulp fibers (CMF) and carboxymethyl-cellulose (CMC) by Sim et al. [26]. These films had a tensile strength of up to 165 MPa and Young’s modulus

196 Advanced Green Composites of 13 GPa. The oxygen permeability was 0.0161 mL m m 2 day 1 kPa 1 at 0% RH. Carboxymethylation of cellulose is a common chemical treatment to facilitate nanofibrillation. However, in this case CMF was used as microsized fibers. The CMC used as matrix/resin is non-toxic and hydrophilic and is even edible. The CMF had portions that were highly swollen which easily collapsed by capillary forces when dried from suspensions in water. The obtained film density was therefore close to that of crystalline cellulose, with practically no pores and resulted in high transparency even with the presence of microfibers. All-cellulose composite films of regenerated cellulose reinforced with cellulose nanowhiskers were introduced by Qi et al. [27]. The cellulosic matrix was obtained by a completely green and unique process using a solution of sodium hydroxide and urea as solvent cooled down to 12 °C. Viscose is the prevailing industrial process to produce regenerated cellulose, but it consumes a significant amount of water and energy besides being water and air pollutant. Carbon disulfide is the most common solvent used in viscose process and is highly toxic [28]. Isotropic films obtained by NaOH/urea dissolved cellulose with 5 wt% whisker content had reasonable strength and an optical transmittance of over 80% at 800 nm. The transparency, though, decreased significantly at higher whisker contents due to easy agglomeration. Yang et al. found that the regenerated cellulose films had very low oxygen transmission rates due to the high crystallinity, with abundant intermolecular hydrogen bonds leading to small free volumes [29]. A film prepared from dissolved cellulose in LiOH/urea/water and coagulated in acetone had an oxygen permeability of 0.003 mL m m 2 day 1 kPa 1, which is an order of magnitude lower than that of cellophane and lower than any plastic film. Films are highly transparent, reasonably bendable, and possess high strength. The drawback of regenerated cellulose films, as well as nanocellulose films is the much higher water vapor permeability than those of polyethylene and polypropylene films. An improved method to produce regenerated cellulose was proposed by Yang et al. [30]. With the addition of zinc oxide to the NaOH/urea solution at 13 °C to dissolve cellulose, regenerated cellulose was coagulated with H2SO4 solution. The obtained regenerated cellulose film had slightly higher optical transparency and tensile strength and modulus of 114 MPa and 6.1 GPa, respectively, compared to corresponding 83 MPa and 3.7 GPa for regenerated cellulose film from NaOH/urea solution. Another eco-friendly way to dissolve cellulose with a 64 wt% H2SO4 aqueous solution was reported by Huang et al. [28]. To restrict the hydrolysis of cellulose chains and allow the cleavage of intermolecular hydrogen bonds, the solution was maintained at 20 °C. The obtained films were

Transparent Green Composites 197 transparent and the mechanical properties were similar to those of films of regenerated cellulose produced using other solvent systems. Yang et al. produced alkali/urea regenerated cellulose films treated by simply soaking them in cationic alkylketene dimer dispersion [31]. After proper drying and heating, these films showed high light transmittance of 88% at 600 nm, tensile strength of 168 MPa, and strain at fracture of 29% with work of fracture of 37 MJ m 3. Films were hydrophobic with a water contact angle of 110° and water uptake of 20%, along with oxygen permeability of 0.13 mL m m 2 day 1 kPa 1 at 50% RH.

8.3 Chitin-Based Transparent Composites 8.3.1 Chitin Nanofiber-Based Composites Chitin, just like cellulose, is another polysaccharide that is found in abundance in nature and is also found in nanofiber form in the hard biocomposite exoskeletons of marine crustaceans, insects and in the cell walls of filamentous fungi. The molecular structures of chitin and cellulose are almost identical, with the only difference that one of the 3 hydroxyl groups in cellulose is replaced by an acetamido group in chitin [32]. While both types of nanofibers have to be extracted from complex biocomposites, chitin nanofibers are relatively easier to obtain than cellulose ones. In the case of chitin, first a chemical treatment is performed to remove proteins and minerals, followed by mechanical shear application [33, 34]. As with cellulose, purified chitin is always kept in a wet state to avoid interfibrillar hydrogen bond formation and aggregation. Ifuku et al. produced chitin nanofibers from crab shell by grinding after chemical purification [35]. The obtained nanofibers had sizes ranging from 10 to 100 nm due to the presence of different structures in the crab shell. Based on an approach proposed by Fan et al., the pH of the purified chitin suspension was decreased to around 3 by adding acetic acid, thereby cationizing the amino groups on the surface of chitin nanofibers to generate electrostatic repulsion [36]. The new protocol delivered nanofibers of 10 to 20 nm in diameter by a single pass through the grinder, with no deacetylation or crystalline structure change. The acidified suspension grinding proved to be effective even with once dried purified chitin, producing the same nanofibers as by the never-dried method. Chitin-reinforced transparent composites were fabricated following the paper sheet impregnation process used in cellulose nanocomposites [37]. Composites formed with acrylic resin with a refractive index 1.532 produced 70 m thick films at a

198 Advanced Green Composites chitin nanofiber loading of 60 wt%. For the chitin nanofiber/acrylic resin composites the regular transmittances at 600 nm wavelength were 89.8%, 88.7%, and 87.9% when never-dried chitin, once-dried chitin and dried powder chitin, respectively, were used as reinforcement. The transmittance decreased less than 2 percentage points as a result of drying the chitin before nanofibrillation. The fibrillation of chitin in acidic medium demands, indeed, less intense mechanical treatment and, as a result, even a kitchen blender was able to produce nanofibers. Shams et al. prepared purified chitin powder from crab shells, which was suspended in an aqueous suspension adding acetic acid to adjust the pH to 3–4 [38]. An agitation by a kitchen blender for 10 minutes produced nanofibers of 20 to 30 nm in diameter. Besides, the time required for filtration to make sheets was 9 times faster than that needed to dewater CNFs of similar size, as an indication of the less hydrophilic character of chitin. The nanocomposites produced by acrylic resin impregnation delivered light transmittances 10 and 20 percentage points higher, at 600 and 400 nm wavelengths, respectively, than those obtained from CNF counterparts. The higher transparency was attributed to the better affinity of chitin with the hydrophobic matrix, for being less hydrophilic than cellulose. The CTE of a 40 wt% loaded chitin nanofiber reinforcement was reduced from 213 × 10 6 K 1 of the acrylic resin to 15.6 × 10 6 K 1, a value comparable to that obtained by CNF reinforcement. The estimated tensile modulus of chitin crystal is 80 GPa [39], almost half the estimated modulus of cellulose crystal at 138 GPa [40]. Considering that the modulus and CTE are both dependent on the atomic potential well, higher modulus would result in low CTE of a crystal. In the case of these composites, again the lower hydrophilicity of chitin compared to cellulose produced better fiber-matrix interactions that might have contributed to the superior restriction of resin expansion. The effect of refractive index mismatch on the optical transparency of chitin nanofiber nanocomposites has also been evaluated [41]. Chitin nanofiber was combined with different acrylic and methacrylic resins with refractive indexes ranging from 1.468 to 1.54. The regular optical transmittances of composites at 600 nm were above 70% at a nanofiber content of about 40%, with the transparency peaking when the resin refractive index was in between 1.50 to 1.51. In some instances the composite transmittance was reduced from that of resin by just 1.3 percentage point. The CTEs were reduced from 98.3 10 6 to 185.4 10 6 K 1 of neat resins to 17.5 10 6 to 29.5 10 6 K 1 of the resulting composites. Silsesquioxane (SSQ) resins are compounds consisting of molecules made of silica cores linked to organo-functional groups that confer them inorganic-organic characteristics. Their inorganic character gives them

Transparent Green Composites 199 higher stiffness and thermal stability than usual polymers, whereas the organic side delivers better processability than conventional ceramics. As a result, SSQs are optically transparent and thermally stable. To counteract the brittleness of SSQs, a bifunctional urethaneacrylate (UA) oligomer was added at varying amounts and copolymerized [42]. At a wavelength of 600 nm the regular transmittance of 50 wt% chitin nanofiber-containing composites was 80 to 85%, while the neat SSQ resin had 90%. The 1:4 ratio SSQ to UA mixture had the transmittance reduced by about 10 percentage points. The CTE of SSQ-UA mixtures ranged from 96 10 6 to 164 10 6 K 1, but chitin nanofiber reinforcement reduced them by 70 to 80% to about 30 10 6 K 1. As expected, decomposition temperatures of composites were higher than chitin alone and increased with increasing SSQ content which is thermally more stable. As nanofiber paper-like sheet impregnation with resin makes it difficult to mold the composites into desired shapes other than flat films, a protocol based on a Pickering emulsion of acrylic resin dispersed in water and stabilized by chitin nanofiber-network was proposed by Shams and Yano [43]. In this process, composite sheets can be obtained by filtration of the Pickering emulsion and posterior drying, followed by hot pressing into a semi spherical shape and UV curing. The compression at the resin’s plasticizing temperature eliminated the nanofiber-resin interfaces and resulted in transparent composites. The regular light transmittance achieved 86.1% with the CTE being reduced to 15.1 10 6 K 1 from 213 10 6 K 1 observed for the neat acrylic resin. These values at a nanofiber content of 14 wt% were practically identical to those of composites obtained by paper-like sheets impregnated with the same resin, but at a higher chitin nanofiber content of 35 to 40 wt%.

8.3.2 Micro-Sized Chitin Composites An interesting demonstration of the presence of chitin in the form of nanofibers in the exoskeleton of crustaceans was reported by Shams et al. [44]. A crab carapace had its proteins, minerals and pigments carefully removed while keeping the original shape made solely by chitin. The interstitial cavities left by the removal of impurities (proteins, minerals and pigments) were filled with acrylic resin, turning the carapace transparent (Figure 8.3). Similarly, transparent composites were also obtained by resin reinforced with chitin powder 0.3 mm in average size, without any fibrillation. The regular transmittance was reduced by just 10 percentage points from that of the neat resin but the CTE decreased from 213 10 6 to just 25 10 6 K 1.

200 Advanced Green Composites

50 mm (a)

(b)

(c)

Figure 8.3 Crab carapaces before chemical treatment (a), after chemical treatment (removal of proteins, minerals and pigments) (b), and after chemical treatment and resin impregnation (c). [44]

Shams and Yano produced composites of acrylic resin reinforced with chitin with differing degrees of fibrillation from micro-sized through sub-micron-sized particles to nanofibers of 30 to 110 nm in diameter [45]. Fibrillation was accomplished in acidic aqueous suspension of purified chitin by a kitchen blender at varying rotational speeds from 1,000 to 37,000 rpm. They found that the fibrils with differing sizes produced over 85% transmittance at 600 nm, as expected. However, even the least fibrillated chitin delivered regular transmittance of above 80%. These values are just 10 and 5 percentage points below the transmittance of neat acrylic resin and nanofiber-reinforced composites, respectively. Meanwhile the CTE of these composites, regardless of fibrillation, were around 18 10 6 K 1 in the temperature range of 20 to 150 °C. Moreover, the chitin-reinforced composites showed lower transmittance decay rates than cellulose nanocomposites during thermal aging consisting of long time exposure to 200 °C, especially at a wavelength of 400 nm. CNFs begin to degrade at around 200 °C due to the less thermally stable hemicelluloses remaining as residues, whereas chitin nanofibers start to degrade at around 280 °C. This shows the advantages of chitin over cellulose as they provide high transparency without the need to nanofibrillate, with consequent faster dewatering times required for papermaking, and higher thermal stability.

8.3.3 Chitin-Chitosan Transparent Green Composites A completely green transparent nanocomposite based on chitin was developed by Ifuku et al. who used partially deacetylated chitin nanofibers to reinforce chitosan resin [46]. Only the surface of chitin nanofibers

Transparent Green Composites 201 were deacetylated in an alkali NaOH solution according to a protocol described by Fan et al. [47] and subsequently nanofibrillated using a grinder. The surface-deacetylated chitin nanofibers were mixed with chitosan dissolved in an aqueous solution of acetic acid and were cast into films. The chitin-chitosan composite films obtained exhibited regular light transmittance of 84% at 600 nm wavelength, which was more transparent than films prepared with surface-deacetylated chitin nanofibers alone at 79.8% [48]. Since chitosan is in fact derived from chitin through deacetylation, the chitin nanofibers are essentially coated with chitosan that makes seamless contact with the matrix, eliminating any interfaces that would scatter light. Increasing the ratio of chitin nanofibers content in the composites did not reduce transparency appreciably; however, the Young’s modulus and tensile strength of the films increased significantly from 2.6 GPa and 40 MPa for neat chitosan to 8.8 GPa and 157 MPa for surface-deacetylated chitin nanofibers, respectively. The CTE decreased from 35.3 10 6 K 1 to 13.5 10 6 K 1, with values below 20 10 6 K 1 for nanofiber contents of 40 wt% and above. In addition, presence of chitosan in the films added anti-bacterial properties to the material as well. The surface-deacetylated chitin nanofibers had glycerol, a plasticizer, mixed in various proportions to modify the mechanical properties of the produced films [49]. As expected, the Young’s modulus and tensile strength of the original nanofiber films decreased to 0.8 GPa and 22 MPa, respectively, at a glycerol content of 70 wt%, but became much more ductile compared to the brittle nanofiber film. An interesting characteristic was that all composite films showed CTE as low as 15 10 6 K 1 regardless of the glycerol (CTE of 490 10 6 K 1) content, due presumably to the formation of a rigid percolated network of chitin nanofibers restricting thermal expansion. A highly transparent chitosan reinforced with CNFs was reported by Fernandes et al. [50]. Films were prepared by casting the mixture, and optical transmittances of 5 wt% nanofiber containing composites were observed to be slightly lower than the transmittance of neat chitosan. Mechanical properties increased consistently with the nanofiber content, and depending on the chitosan resin type the tensile modulus increased by 320% over that of pure matrix. Chitosan reinforced with shredded BC pellicle was also developed by Fernandes et al. [51]. As with CNF reinforcement, the light transmittance in the 400–700 nm range of pure chitosan of about 90% barely decreased with addition of 5 wt% BC, however 10 wt% nanofiber content reduced transmittance to 80%. There was also some indication of low oxygen permeability of films with BC-reinforced chitosan films.

202 Advanced Green Composites

8.3.4 All Chitin Nanofiber Transparent Films Another approach to extract chitin nanofibers was by esterification of their surface hydroxyl groups with maleic anhydride and subsequent nanofibrillation by a grinder [52]. Anionic charges on the surface of nanofibers produce electrostatic repulsion and osmotic pressure which facilitates individualization of nanofibers. The average diameter of the nanofibers obtained was about 11 nm while the diameter of those from the unmodified chitin averaged 45 nm. After casting into films, the transparency of maleated chitin nanofiber film at 600 nm wavelength was 74.4%, whereas that of unmodified chitin nanofiber was only 2.4%. This is another example of a fully green transparent composite made solely of chitin nanofibers modified by a compound of low environmental hazard.

8.4

Electronic Devices Based on CNF Films and Composites

The use of CNF-based transparent composite films as substrates for flexible electronics has been extensively studied [53–59]. Nogi et al. developed foldable solar cells using nanopapers as substrates [53]. The simple protocol consisted of depositing a silver nanowire suspension on the surface of a silicon wafer and subsequently dropping a CNF suspension on the nanowire layer. As the dropped suspension shrunk only in thickness, the capillary forces involved in the drying of CNFs acted as a self-pressing mechanism to assure physical contact between the Ag nanowires. The paper film peeled from the substrate had sheet resistance of 17 sq 1 and an optical transmittance of 94.4% at 600 nm wavelength, values as high as those of indium tin oxide (ITO). The conductive nanopaper was flexible, foldable and could keep its conductivity after 20 folding cycles. Nanopaper solar cell manufactured via printing of organic components presented a power conversion efficiency of 3.2%, comparable to those of ITO-based counterparts, but keeping its conversion efficiency during and after folding. Koga et al. developed electrically conductive composite films by combining carbon nanotubes (CNT) as reinforcing agent and TOCN as the resin [54]. As they noticed that CNT could be dispersed to a certain degree in water in the presence of TOCN, CNT was acid-treated to attach carboxyl groups to their surfaces so that electrostatic repulsion could further improve dispersibility in water. The TOCN and CNT aqueous suspensions were mixed to produce thin translucent films, by casting, of

Transparent Green Composites 203 around 5  m in thickness with optical transmittance of about 30% at 600 nm wavelength. The electrical conductivity was superior compared to those of CNT-filled polymers such as regenerated cellulose, BC, polyethylene, PS, and polyamide-6, and was sufficient for using in touch-panel applications. In addition, the authors showed the possibility of ink-jet printability of patterns on substrates, by casting CNT-TOCN nanoink on polyethylene terephthalate (PET) films. The sheet resistance was 1.2 k sq 1 and light transmittance dropped from 80% of PET to 70% after deposition, showing the effectiveness of TOCN as a dispersing agent of CNT. As a substitute for brittle ITO, as a transparent and electrically conductive material, an alternative transparent, conductive, and foldable material was fabricated by deposition of silver nanowires or CNTs on cellulose nanopaper [55]. In this case aqueous suspensions of CNFs, 15 nm in diameter, were dewatered by vacuum filtration and afterwards another aqueous suspension of silver nanowires was vacuum filtered on top of the wet CNF sheet. The wet cellulose and silver nanowire layers were then hot-pressed under slight pressure at 110 °C for drying. This silver nanowire-coated CNF paper had sheet resistance of 12 sq 1, with optical transmittance of 88% at 550 nm wavelength and haze of just over 12%. The sheet resistance was 75 times lower than that of PET film substrate covered by silver nanowires by drop coating process and improved the performance of films obtained by conventional coating techniques. The silver nanowires formed a randomly oriented and percolated network by filtration, and the compression during drying embedded it into the CNF substrate producing strong adhesion. The robustness of the material was demonstrated by multiple foldings of the sheet without any change in electrical resistance. Using the same nanopaper as a substrate, Fujisaki et al. deposited an organic thin-film transistor (OTFT) array on it using lithographic and solution-based processes suited for large-area and lowcost manufacturing [56]. The obtained device had high thermal stability of over 180 °C, excellent resistance to solvents and low CTE. Due to the smoothness of the nanopaper surface, the OTFT device had high mobility and was almost free of hysteresis, with good operational and mechanical stabilities. Unlike CNFs obtained by TEMPO-mediated oxidation (TOCN), mechanically fibrillated nanofibers kept high light transmittance at 600 nm of 85.4%, slightly down from 86.7%, when subjected to 180 °C for 1 hour in air. However, Yagyu et al. were successful in making transparent TOCN papers more resistant to yellowing, due to heat, by decreasing the carboxylate content of chemically modified cellulose during extraction treatment [57].

204 Advanced Green Composites Another example of flexible electronics is a flexible nonvolatile memory produced by growing silver nanoparticles in-situ on the surface of CNFs extracted from wood by TEMPO-mediated oxidation [58]. The flexible memory had a nonvolatile resistive switching with 6 orders of ON/OFF resistance ratio and small variation in operation voltage. It had exceptional flexibility without performance degradation up to a bending radius of 350 m, at the time it was the smallest value reported for this kind of device. Perhaps the most sophisticated application of CNFs to prepare transparent substrates for electronic devices was the one reported by Ji et al. [59]. At a wavelength of 550 nm, epoxy films had an optical transmittance of 90% and haze of 3% whereas CNF films had a lower transmittance of 77% and haze of 18% due to cavities and surface roughness that scatter light. A combination of both materials was realized by a reinforcing hybrid backbone made by TOCN suspension being simultaneously sprayed on epoxy nanofibers during formation by electrospinning. This TOCN/ epoxy nanoweb was then embedded in an epoxy resin as indicated in the diagram of the composite film fabrication shown in Figure 8.4. The optical transmittance of the obtained hybrid film was 88% with 4% haze at 550 nm wavelength. The tensile strength and modulus were 220 MPa and 6 GPa for TOCN films and 40 MPa and 1.7 GPa for pristine epoxy, respectively. Hybrid epoxy films exhibited strength of 190 MPa and modulus of 5.5

CNF spraying

Epoxy electrospinning

(b)

(c)

(a)

(d)

Figure 8.4 SPreparation of the CNF hybrid film. (a) A schematic image of an electrospinning epoxy backbone and spraying CNF fillers simultaneously. Schematics and scanning electron microscopy (SEM) images (top and cross-sectional views) of (b) the electrospun three-dimensional (3D) epoxy nanoweb structure, (c) the CNF–epoxy hybrid before hot pressing and (d) the CNF–epoxy hybrid after hot pressing. White scale bars are 10 μm, and black scale bars are 20 μm. [59].

Transparent Green Composites 205 GPa. The CTE of epoxy at 53 × 10 6 K 1 was reduced to 10 × 10 6 K 1. A flexible touchscreen panel was fabricated on the hybrid film by depositing an silver nanowire electrode that reduced light transmittance to 83% at 550 nm wavelength. The electrode’s initial sheet resistance of 12 sq 1 was barely changed at a bending radius of 270 m. Cyclic bending at a radius of 0.5  mm and frequency of 0.5 Hz increased the sheet resistance of the sheet by 11% for tensile and 5% for compressive bending after 10,000 cycles. Next, polymer dot spacers and gold interconnects were patterned on the hybrid epoxy film by photolithography and metal evaporation. The device was shown to be operationally stable for 10 days under exposure to a temperature of 100 °C. A transparent organic light-emitting diode was also fabricated on the hybrid film, showing a light transmittance of 62% at 550 nm. The maximum luminance of the panel was 410.3 cd m 2, which was comparable to that fabricated on glass substrate. The authors even proposed the manufacture of photo-patternable origami substrates with reversible fold-ability and stretch-ability through elastomeric joints using these hybrid composite films, allowing the manufacture of devices with complex shapes.

8.5 Future Prospects Even though the first transparent composites based on cellulose or chitin were not fully ‘green’, they soon evolved to transparent paper-like films and all-cellulose composites. The use of biodegradable resins also turned composites into fully green materials. As many of these films have low oxygen permeability, they seem to become the obvious choice for food packaging, a low tech application but with a vast worldwide market. However, the development of transparent composites aimed at electronic devices is still on the way to become fully green. The need for a smoother surface for device layer deposition and reduced haze requires the use of a polymeric matrix that is still petroleum-based. However, with the development of new bioplastics such as lignin-based epoxies, green alternatives to petroleum-based polymers would soon become available. But bio-based reinforcing nanofibers are essential for these composites as the best and perhaps the only way to restrict CTE and keep transparency to desired levels and stay green. As long as displays’ backplane technologies are based on brittle inorganic compounds, the CTEs of substrates have to be constrained to match those of the inorganic compounds used. At least until the day the organic electronics, which are more ductile, become commonplace.

206 Advanced Green Composites

8.6 Summary The semi-green transparent composites that started with cellulose and chitin nanofibers reinforcing synthetic polymers have rapidly evolved into fully green materials as nanofiber-based paper films and all-cellulose or all-chitin composites. While the high transparency has been maintained and the thermal expansion constrained to values adequate for applications in electronic devices, they can now be made from totally biobased materials possessing easy disposability which results from their inherent biodegradability. A new property has been identified in some of these all-cellulose films, which is their low oxygen permeability as a result of the combination of a massive number of interfibrillar hydrogen bonds and tiny non-connected voids of tens of nanometers in size. This characteristic makes these transparent films suitable for food packaging, with a huge market demand but has a desperate need for biodegradable alternatives to replace environmentally harmful conventional plastics.

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208 Advanced Green Composites 23. H. Fukuzumi, T. Saito, S. Iwamoto, Y. Kumamoto, T. Ohdaira, R. Suzuki and A. Isogai, Pore Size Determination of TEMPO-Oxidized Cellulose Nanofibril Films by Positron Annihilation Lifetime Spectroscopy, Biomacromolecules 12, 4057–4062, 2011. 24. X. Sun, Q. Wu, S. Ren and T. Lei, Comparison of highly transparent all-cellulose nanopaper prepared using sulfuric acid and TEMPO-mediated oxidation methods, Cellulose 22, 1123–1133, 2015. 25. H. Fukuzumi, T. Saito, T. Iwata, Y. Kumamoto and A. Isogai, Transparent and High Gas Barrier Films of Cellulose Nanofibers Prepared by TEMPOMediated Oxidation, Biomacromolecules 10, 162–165, 2009. 26. G. Sim, Y. Liu and T. van de Ven, Transparent composite films prepared from chemically modified cellulose fibers, Cellulose 23, 2011–2024, 2016. 27. H. Qi, J. Cai, L. Zhang and S. Kuga, Properties of Films Composed of Cellulose Nanowhiskers and a Cellulose Matrix Regenerated from Alkali/Urea Solution, Biomacromolecules 10, 1597–1602, 2009. 28. W. Huang, Y. Wang, L. Zhang and L. Chen, Rapid dissolution of spruce cellulose in H2SO4 aqueous solution at low temperature, Cellulose 23, 3463–3473, 2016. 29. Q. Yang, H. Fukuzumi, T. Saito, A. Isogai and L. Zhang, Transparent Cellulose Films with High Gas Barrier Properties Fabricated from Aqueous Alkali/Urea Solutions. Biomacromolecules 12, 2766–2771, 2011. 30. Q.  Yang, X.  Qin and L.  Zhang, Properties of cellulose films prepared from NaOH/urea/zincate aqueous solution at low temperature, Cellulose 18, 681– 688, 2011. 31. Q. Yang, T. Saito and A. Isogai, Facile fabrication of transparent cellulose films with high water repellency and gas barrier properties, Cellulose 19, 1913– 1921, 2012. 32. E. S. Stevens, Green plastics: An introduction to the new science of biodegradable plastics, first edition, pp. 83–103, Princeton University Press, 2002. 33. K. G. Nair and A. Dufresne, Crab shell chitin whisker reinforced natural rubber nanocomposites. 1. Processing and swelling behavior, Biomacromolecules 4, 657–665, 2003. 34. J. N. BeMiller and R. L. Whistler, Alkaline Degradation of Amino Sugars, J. Org. Chem. 27, 1161–1164, 1962. 35. S. Ifuku, M. Nogi, K. Abe, M. Yoshioka, M. Morimoto, H. Saimoto and H. Yano, Preparation of Chitin Nanofibers with a Uniform Width as alphaChitin from Crab Shells, Biomacromolecules 10, 1584–1588, 2009. 36. Y.  Fan, T. Saito and A. Isogai, Preparation of chitin nanofibers from squid pen beta-chitin by simple mechanical treatment under acid conditions, Biomacromolecules 9, 1919–1923, 2008. 37. S. Ifuku, M. Nogi, M. Yoshioka, M. Morimoto, H. Yano and H. Saimoto, Fibrillation of dried chitin into 10–20 nm nanofibers by a simple grinding method under acidic conditions, Carbohyd. Polym. 81, 134–139, 2010.

Transparent Green Composites 209 38. M. I. Shams, S. Ifuku, M. Nogi, T. Oku and H. Yano, Fabrication of optically transparent chitin nanocomposites, Appl. Phys. A-Mater. 102, 325–331, 2011. 39. M. Wada and Y. Saito, Lateral thermal expansion of chitin crystals, J. Polym. Sci. Pol. Phys. 39, 168–174, 2001. 40. T. Nishino, K. Takano and K. Nakamae, Elastic-Modulus of the Crystalline Regions of Cellulose Polymorphs, J. Polym. Sci. Pol. Phys. 33, 1647–1651, 1995. 41. S. Ifuku, S. Morooka, A. N. Nakagaito, M. Morimoto and H. Saimoto, Preparation and characterization of optically transparent chitin nanofiber/ (meth)acrylic resin composites, Green Chem. 13, 1708–1711, 2011. 42. S. Ifuku, A. Ikuta, T. Hosomi, S. Kanaya, Z. Shervani, M. Morimoto and H. Saimoto, Preparation of polysilsesquioxane-urethaneacrylate copolymer film reinforced with chitin nanofibers, Carbohyd. Polym. 89, 865–869, 2012. 43. M. I. Shams and H. Yano, Doubly curved nanofiber-reinforced optically transparent composites, Sci. Rep. - UK 5, 16421, 2015. 44. M. I. Shams, M. Nogi, L. A. Berglund and H. Yano, The transparent crab: Preparation and nanostructural implications for bioinspired optically transparent nanocomposites, Soft Matter 8, 1369–1373, 2012. 45. M. I. Shams and H. Yano, Simplified Fabrication of Optically Transparent Composites Reinforced with Nanostructured Chitin, J. Polym. Environ. 21, 937–943, 2013. 46. S. Ifuku, A. Ikuta, M. Egusa, H. Kaminaka, H. Izawa, M. Morimoto and H. Saimoto, Preparation of high-strength transparent chitosan film reinforced with surface-deacetylated chitin nanofibers, Carbohyd. Polym. 98, 1198–1202, 2013. 47. Y. Fan, T. Saito and A. Isogai, Individual chitin nano-whiskers prepared from partially deacetylated alpha-chitin by fibril surface cationization– Carbohyd. Polym. 79, 1046–1051, 2010. 48. Y.  Fan, H. Fukuzumi, T. Saito and A. Isogai, Comparative characterization of aqueous dispersions and cast films of different chitin nanowhiskers/ nanofibers, Int. J. Biol. Macromol. 50, 69–76, 2012. 49. S. Ifuku, A. Ikuta, H. Izawa, M. Morimoto and H. Saimoto, Control of mechanical properties of chitin nanofiber film using glycerol without losing its characteristics, Carbohyd. Polym. 101, 714–717, 2014. 50. S. C. M. Fernandes, C. S. R. Freire, A. J. D. Silvestre, C. P. Neto, A. Gandini, L. A. Berglund and L. Salmen, Transparent chitosan films reinforced with a high content of nanofibrillated cellulose, Carbohyd. Polym. 81, 394–401, 2010. 51. S. C. M. Fernandes, L. Oliveira, C. S. R. Freire, A. J. D. Silvestre, C. P. Neto, A. Gandini and J. Desbrieres, Novel transparent nanocomposite films based on chitosan and bacterial cellulose, Green Chem. 11, 2023–2029, 2009. 52. Y. F. Aklog, T. Nagae, H. Izawa, M. Morimoto, H. Saimoto and S. Ifuku, Preparation of chitin nanofibers by surface esterification of chitin with maleic anhydride and mechanical treatment, Carbohyd. Polym. 153, 55–59, 2016.

210 Advanced Green Composites 53. M. Nogi, M. Karakawa, N. Komoda, H. Yagyu and T. T. Nge, Transparent Conductive Nanofiber Paper for Foldable Solar Cells, Sci. Rep. - UK 5, 17254, 2015. 54. H. Koga, T. Saito, T. Kitaoka, M. Nogi, K. Suganuma and A. Isogai, Transparent, Conductive, and Printable Composites Consisting of TEMPO-Oxidized Nanocellulose and Carbon Nanotube, Biomacromolecules 14, 1160–1165, 2013. 55. H. Koga, M. Nogi, N. Komoda, T. T. Nge, T. Sugahara and K. Suganuma, Uniformly connected conductive networks on cellulose nanofiber paper for transparent paper electronics, NPG Asia Mater. 6, e93, 2014. 56. Y. Fujisaki, H. Koga, Y. Nakajima, M. Nakata, H. Tsuji, T. Yamamoto, T. Kurita, M. Nogi and N. Shimidzu, Transparent Nanopaper-Based Flexible Organic Thin-Film Transistor Array, Adv. Funct. Mater. 24, 1657–1663, 2014. 57. H. Yagyu, T. Saito, A. Isogai, H. Koga and M. Nogi, Chemical Modification of Cellulose Nanofibers for the Production of Highly Thermal Resistant and Optically Transparent Nanopaper for Paper Devices, ACS Appl. Mater. Inter. 7, 22012–22017, 2015. 58. K. Nagashima, H. Koga, U. Celano, F. Zhuge, M. Kanai, S. Rahong, G. Meng, Y. He, J. De Boeck, M. Jurczak, W. Vandervorst,T. Kitaoka,  M. Nogi and T. Yanagida, Cellulose Nanofiber Paper as an Ultra Flexible Nonvolatile Memory, Sci. Rep. - UK 4, 5532, 2014. 59. S. Ji, B. G. Hyun, K. Kim, S. Y. Lee, S.-H. Kim, J.-Y. Kim, M. H. Song and J.-U. Park, Photo-patternable and transparent films using cellulose nanofibers for stretchable origami electronics, NPG Asia Mater. 8, e299, 2016.

9 Toughened Green Composites: Improving Impact Properties Koichi Goda Dept. of Mechanical Engineering, Yamaguchi University, Ube, Yamaguchi, Japan

Abstract Green composites are being used in many applications. However, for them to increase their presence in even more applications, a key research area would be needed to improve their toughness. In the present chapter, toughness of green composites is discussed, especially from the viewpoint of improving their impact properties. In the first section, the effect of fiber length on the stiffness, static strength and impact properties of non-green composites, such as glass fiber composites have been introduced to obtain general understanding of the impact properties of normal fiber-reinforced composites. In the second section, interfacial strength between natural fibers and resin in green composites has been described as a key point for improving the impact properties. Fiber length in a resin is also mentioned as it is deeply related with the degree of impact strength. In general, static strength and stiffness are improved by compounding natural fibers into resin, while the impact strength is reduced. On the other hand, high impact properties are a result of the toughened resin, while the strength and stiffness properties are reduced at the same time. The third section describes a solution for such conflicting properties, how toughened natural fibers could be used for green composites as a way to improve their toughness. In the last section, the fiber structure and mechanisms to increase their toughness have been described. Keywords: Green composites, natural fibers, thermoplastic resin, strength, stiffness, toughness, impact properties

9.1 Introduction Improving the toughness of green composites is one of the most important needs if these materials are to be used in a wide range of industrial products. Corresponding author: [email protected] Anil Netravali (ed.) Advanced Green Composites, (211–246) © 2018 Scrivener Publishing LLC

211

212 Advanced Green Composites In general, the entire area under the stress/strain curve, up to the failure of the specimen (resin, fiber or composite), is used as the measure of the work of fracture or modulus of toughness. Impact resistance or the energy absorbed during impact, an index of materials tested under impact loads, has often been related with the measure of toughness. In this chapter, thus, impact properties of green composites reinforced with natural fibers are discussed as the main topic. Especially, the focus is on what role natural fibers can play in order to toughen relatively low toughness conventional polymers such as polypropylene (PP) as well as ‘green’ polymers such as polylactic acid (PLA).

9.2 Significance of Fiber Length in Toughened Fibrous Composites To begin with, it would be beneficial to briefly review the mechanisms to obtain an increase in the impact strength of synthetic, petroleum-based fiber-reinforced composites, which have been studied for several decades. It is commonly known that stiffness, static tensile strength and impact strength of glass fiber reinforced polymer composites depend on fiber length as well as fiber/resin interfacial strength, because the degree of stress transfer from matrix to fibers changes depending on the length of the fiber. Furthermore, the optimal fiber lengths for enhancing stiffness, static tensile strength or impact strength are different. Figure 9.1 shows a schematic of the relation between mechanical properties and fiber length, of which the vertical axis has been normalized. As seen in Figure  9.1, modulus (stiffness) of the composite increases rapidly even if the fiber lengths are quite small, and achieves the maximum value, while the impact strength (Charpy) increases gradually, in comparison. When the fiber lengths correspond to the saturating level of stiffness, the impact strength remains at a low level. Static tensile strength is placed in the middle, between these two curves. Thus, it is understood that the fiber length is an important structural parameter in deciding mechanical properties and must be large for toughening fibrous composites. This also implies that impact strength is structure-sensitive, while stiffness is structure-insensitive. To discuss quantitatively, the change in mechanical properties in Figure  9.1, utilizing some theoretical models would be useful. For example, Cox model [1] is a representative theory to predict stiffness or tensile modulus (Young’s modulus), Ec, of short fiber-reinforced composites. When the rule of mixture is used, this model can be given as follows:

Ec

Vf E f

0 l

(1 V f )Em

(9.1)

Toughened Green Composites: Improving Impact Properties 213 1

Normalized value

0.8

0.6

0.4

Strength Modulus Charpy impact

0.2

0 2

4 6 8 Fiber length (mm)

10

12

Figure 9.1 Influence of fiber length on normalized glass fiber-reinforced composites (Data plots in Figure 16 [4] are converted into numerals by Simple Digitizer, and plotted again in this graph)

where, Ef and Em are Young’s moduli of fiber and matrix resin, respectively. Vf is the fiber volume fraction, 0 is a fiber orientation factor, and l is the fiber length efficiency factor, given as:

l

1

tanh( l / 2) l /2

(9.2)

where, 1/2

2 Em 1 r E f (1 m )ln(r / R)

,

l is the fiber length, r is the fiber radius, and R is the distance between centers of fibers. m is a Poisson’s ratio of the matrix. Further, the term ln (r/R) is often expressed using the fibers’ geometrical arrangement, Pf, as

ln

r R

Pf 1 ln 2 Vf

(9.3)

where, Pf = for square pitch, and Pf 2 / 3 for hexagonal array [2]. As shown in the above equation, l is a factor that decides the effective Young’s modulus of short fiber-reinforced composites. Mathematically, l

214 Advanced Green Composites rises quickly with increase in l and converges to a unity. The shape of this behavior is a convex upward, which resembles the change in tensile modulus shown in Figure 9.1. On the other hand, Kelly-Tyson model, based on matrix plasticity at the fiber surface near the fiber ends, has often been used for strength prediction of short fiber-reinforced composites [3]. The rule of mixture in eq. (9.1) is changed, as a function of fiber strength f, as follows: c

Vf

0 l

f

(1 V f )

m

(9.4)

where, c is the strength of the composite and ’m is matrix stress corresponding to fiber fracture strain. In this model, l is given as:

(1 lc / 2l )

l

(9.5)

where lc is a length deciding the reinforcing efficiency, called the critical fiber length. In other words, to utilize the fiber strength, the fiber length must be larger than lc (often called ‘supercritical’ fibers). The function mentioned above also rises with increase in l, and shows an upward convex. In the case where lc > l (i.e. ‘subcritical’ fibers), the l factor is given as: l

l / 2lc

(9.6)

This is a linear function of l, and therefore the degree of rise may become slower than obtained from the form given in eq. (9.2). These two functions given in eqs. (9.5) and (9.6), are switched (numerator to denominator) with respect to the length lc, and, therefore, the saturation level of composite strength with increase in l would be delayed as compared to tensile modulus. In other words, the behaviors of tensile modulus and tensile strength with increase in l shown in Figure 9.1 are supported from the mathematical viewpoint. Thomason et al. [4], on the other hand, reported that Cottrell model showing impact energy is not in agreement with the trend of experimental data. This model shows a fiber pull-out dominant model, the concept which has been discussed by many researchers including Kelly [5] and Cooper [6]. According to the Cottrell model, the energy initially increases with a downward convex, but decreases with a downward convex after l exceeds lc. Actually, the calculation by Kelly and Cooper showed that the fiber pullout energy changes in proportion to l2 (See, Appendix). Matthews and Rawlings [7] also pointed out that the fiber pull-out energy is much higher than the debonding energy from their calculation. As a result, they stated

Toughened Green Composites: Improving Impact Properties 215 that the fiber pull-out from resin is more significant toughening mechanism in fiber reinforced composites. As shown in Figure 9.1, however, the experimental data shifts with an upward convex. This is quite opposite to the behavior with proportion to l2. Thomason et al. [4] pointed out for their long glass fiber composites (randomly arrayed in-plane) that fracturing of fibers was an important mechanism in the energy absorption process from a positive correlation between tensile and impact strengths. They also suggested that the impact energy changes in proportion to the ratio l/(l+lc), which can give an upward-convex behavior with a quite slow convergence to unity, as shown in Figure  9.1. The same behavior has also been confirmed for flax fiber-reinforced composites [8]. On the other hand, if the composite was reinforced with the subcritical fibers, i.e., shorter than lc, the energy absorption is mainly dominated by pull-out mechanism [9]. Norman and Robertson calculated energies of fiber pull-out, fiber fracture, matrix yielding and fiber–matrix debonding during the composite fracture process, and discussed which energy is the most influential [9]. Their results show that, snubbing is predicted to provide a significant contribution to pull-out toughening. Fiber breakage is less effective than fiber pull-out because only the broken ends pull-out, and also that it occurs for fibers with > c ( is the fiber angle between the fiber axis and the normal to the fracture plane and c is the critical fiber angle, which is 7–8°). Thus, the dominant mechanism of fracture energy by impact test is largely dependent on fiber length. That is also to say that the fracture energy of long fiber (supercritical fiber) composites is fiber-fracture dominant, while that of short fiber (subcritical fiber) composites is dominated by pull-out mechanism. A schematic of the fracture mechanism for the two composites is shown in Figure 9.2. In addition to both fiber fracture and pull-out, matrix plasticity also contributes to the fracture energy. Lauke et al. [10] have considered the fracture process in short glass fiber-reinforced composites containing fibers of subcritical length. Their micro-mechanical model suggests that the total work of fracture contains contributions from debonding, sliding or pull-out, and plastic deformation of the matrix. Under static loading conditions, intensive plastic flow occurs locally in the matrix and the contribution of the matrix to the work of fracture in this case is predominant. Under dynamic loading (impact tests), the matrix responds in a more brittle way and the debonding and pull-out contributions tend to dominate. Gupta et al. [11] also accepted this mechanism through their experimental results. To exhibit the plastic deformation of matrix more largely under dynamic loading, plasticization is carried out for the matrix (resin). To do this, liquid (low molecular weight) polymers have been used as plasticizers

216 Advanced Green Composites Crack path

Fiber-fracture dominant Fracture energy

Pull-out dominant

Contribution of pull-out energy

lc

Fiber length l

Figure 9.2 A schematic of fracture energy as a function of fiber length in a fiberreinforced composite under impact force.

from old times [12]. The plasticization treatment is often carried out, in the case of thermoset resins such as epoxy or unsaturated polyester [13]. They reported that by containing 15 phr acrylonitrile butadiene rubber into unsaturated polyester resin, impact strength (Izod, unnotched) of bulk molding compound (BMC) based composites increased from 1.25 J/ cm to 1.64 J/cm, over 30% increase. However, the addition of such rubbers is often accompanied by decrease in the corresponding tensile strength and Young’s modulus values. Tensile strength and modulus of the BMC composites decreased to 90% level as compared to the original composites without the rubber additive. In recent years, PP has seen a gradual increase in its use in composites as a matrix material, especially when injection molding has been used for production. In these types of composites as well, improvement in impact strength has been carried out. One of the effective methods for increasing impact strength of glass fiber/polypropylene (GF/PP) composites is by changing crystalline structure of the PP. -crystals, one of PP crystal structures, has received the most attention due to the excellent impact toughness and elongation at break [14, 15]. The commonly obtained -crystals offer excellent stiffness but inferior toughness because of the interlocking effect of the radial lamellae by the tangential crystallites that make plastic deformation difficult. In contrast, -crystals without the cross-hatching show

Toughened Green Composites: Improving Impact Properties 217 lower stiffness than -crystals but higher toughness. Modification through -crystallization improves the impact strength (Charpy, notched) of GF/ PP without significant decrease in tensile strength [16]. Furthermore, Geng et al. confirmed the increase in impact strength (Izod, notched) at a wide range of temperatures and glass fiber contents without significant loss of strength and modulus [17]. Another efficient way to enhance the toughness of PP matrix is by annealing them. It was demonstrated, using isotactic polypropylene (iPP), that the highest impact strength (Izod, notched) can be obtained after annealing it at 130 C, as compared to the untreated specimens, although their tensile strength decreased slightly [18]. Impact strength (Izod, notched) of PP containing a small amount of polyethylene (PE) oxide was also enhanced after annealing at 130–140 C [19]. In the latter paper, additionally, areas under the stress-strain curves in tensile tests were measured and were shown to have the highest values after annealing at 140 C. Through such annealing treatments, it was confirmed that impact strengths of GF/PP composites did improve [17]. In addition, they reported that, by combining -crystallization with annealing, impact strength increased approximately five times compared to neat PP resin at room temperature. Thus, the field of glass fiber-reinforced plastic composites has been progressing toward a highly toughened material.

9.3 Impact Properties of Green Composites 9.3.1 Relation Between Interfacial and Mechanical Properties in Green Composites In general, static strengths, such as tensile and bending strengths, of fibrous composites increase with increase in fiber content, irrespective of the difference between natural and synthetic fibers. Stiffness or Young’s modulus also increases with increase in the fiber content. However, some green composites have not achieved such improvements in their properties that accompany the increase in fiber content. According to Eng et al. [20], PLA matrix based composites reinforced with oil palm mesocarp fibers (OPMF) decreased in strength and stiffness as well as impact strength (Izod, notched), as shown in Table 9.1. If a silane-coupling agent is used for the fibers to improve the fiber/resin interfacial properties, the degree of decrease in strength and stiffness can be reduced. However, these properties were still less than those obtained for neat PLA. This means, although the coupling treatment was somewhat effective, the interfacial bond between OPMF and PLA was still insufficient. Combination of chicken feather fiber

Izod (notched)

Bagasse

Compression molding

Izod Extrusion + (Unnotched) injection

Chicken feather

Compression molding

Production method

Izod (notched)

Type of impact test

Oil palm mesocarp fiber (OPMF)

Type of fiber

29.5

0 30

6761

6761/RBF 16.7 24.4

6761/PRF

6761/EBF

22.3

21.6 24.4

9856B/EBF

26.2

9856B/PRF

30

9856B/RBF

32.3

36.8

10 0

47.4 41.4

2 5

9856B

Long fiber/PLA

38.6

2.34*

2.32*

2.78*

1.61*

2.77*

2.49*

2.30*

1.86*

3.12

3.13

3.08

3.17

3.08

43.1

5 10

3.00 3.08

2

Short Fiber/PLA

56.7

0.768

0.843

0.866

0.577

0.661

0.686

0.903

47.4

0

29.4

30

PLA

33.2

20

30 40.5

29.5

20 10

32.3

10

PLA/untreated OPMF

PLA/silane-treated OPMF

36.3

0

45.9

Tensile Tensile strength modulus (MPa) (GPa)

PLA

Type of specimen

Fiber content (wt%)

2.37

3.33

3.87 0.608

1.32

1.05

1.05 0.70

3.63

1.31

1.61

2.56

4.13

4.97

1.38

1.40

2.11

1.36

1.42

1.64

3.73

38.2

35.8

32.8

39.7

60.9

45.9

41.9

73.8

94.6

120

142

66.8

78.5

95.8

188

Lu, et al. 2006 [24]

Baba, et al. 2015 [21]

Eng, et al. 2014 [20]

Impact Elongation Impact at break strength strength (J/m) Ref. (%) (kJ/m2)

Table 9.1 Mechanical properties of green composites with poor interfacial bond.

Note

Any aspect ratio of +20# is larger before and after compounding process HDPE resins. RBF: raw bagasse fiber PRF: pure rind fiber EBF: alkali-extracted bagasse fiber *Tensile modulus shows storage modulus E’ Izod impact testing: ASTM D 256

Data plots in Fig. 10 of this reference are converted into numerals by Simple Digitizer. Short fiber: 3mm pellet length Long fiber: 20mm pellet length

Resin: 85%wtPLA+15%wtPCL (PCL: polycaprolactone) 1wt%nanoclay containing in the resin Elongations at break are converted to strain by dividing the orignal data by grip attachment distance of 45mm. Izod impact testing: ASTM D 256

218 Advanced Green Composites

Izod (notched)

Izod (notched)

Izod, unnotched

Wood, rice husk, straw leaf, straw stem, whole straw

Bagasse, wood

Oil palm empty fruit bunch

Single-screw extruder + compression molding

Compression molding

Compression molding

OPEFB/PP(EX) 13.8 13.2 10.7 9.29

30 40 50 60

15

60 14.1

15.5

50

20

18.5

40

13.2

18.4

30

0

18.6

20

36.1

2.4wt%MAPE/pine 18.6

32.6 0

2.23*

23.5

1.2wt%MAPE/pine

OPEFB/PP(IM)

2.50*

29.8

0.419

0.306

0.262

0.258

0.262

0.300

0.891

0.788

0.713

0.538

0.439

0.300

2.26*

2.34*

2.50*

2.29*

1.55*

2.97*

3.31*

2.98*

2.60*

2.51*

2.24*

2.48*

2.13*

1.98*

30.6

21.4

1.55* 2.51*

1.5wt%MAPE/ bagasse 3.0wt%MAPE/ bagasse pine

15

bagasse

32.6

19.4

Whole straw 0

19.3

R-HDPE

19.1

23.1

Wood

Straw stem

21.5

Whole straw

Straw leaf

21.2

Straw stem

17.3

21.7

Rice husk

19.7

50

32.4 23.1

Straw leaf

Wood

Rice husk

0 30

R-HDPE

2.72

4.14

6.21

7.05

7.82

12.3

2

2.59

3.62

4.98

6.01

6.85

6.09

6.74

6.10

5.90

5.70

4.92

12.3

4.24

4.85

3.60

4.37

5.60

6.62

7.56

5.99

7.42

5.60

13.2

185

60.1

56.8

62.1

63.1

68.4

97

74.7

90.3

101

133

149

R-HDPE: Recycled high density polyethylene *Tensile modulus shows storage modulus E’

Rozman, Data plots in Figs. 1-4 of this reference are converted into et al. numerals by Simple Digitizer. 1998 [27] IM: internal mixer EX: singlescrew extruder Izod impact testing: ASTM D 256

Lei, et al. R-HDPE: Recycled high density 2007 polyethylene [26] MAPE: maleated polyethylene

Yao, et al. 2008 [25]

Toughened Green Composites: Improving Impact Properties 219

220 Advanced Green Composites (CFF) with PLA also reduced tensile strength as well as impact strength with increase in CFF content [21]. On the other hand, tensile modulus increased slightly, which means, CFF was stiffer than neat PLA. Poor interfacial bonding properties are not efficient in transferring an external force to the fibers, which leads to low static strength properties [22], as well as impact strength. Therefore, if there is sufficient interfacial bond between CFF and PLA, tensile strength should increase with increase in CFF content, as described in the previous section, but its actual behavior shows an opposite tendency. Bajpai et al. described in their review article of PLAbased green composites that for good impact strength, a most favorable bonding is necessary [23]. PLA is a brittle polymer and shows relatively low elongation at break, but even a ductile polymer such as PE also induces lower tensile properties of composites than the neat resin. Lu et al. produced sugarcane bagassebased fiber composites using four types of high-density PE (HDPE) varieties by compression molding [24]. After fiber addition, tensile strength of the composites decreased to lower than that of neat HDPE, although the storage moduli were higher than the neat resin level except for the weakest HDPE (not shown in Table 9.1). Impact strength (Izod, notched) also decreased, as compared to that of the neat resin. Table 9.1 shows the results of only two HDPE composites with higher strength. The authors implied that the low mechanical properties were a result of the poor interfacial properties that could be possibly resolved by the use of coupling agents. Such lower strength properties lower than the neat resin, have also been reported in green composites using wood, rice husk, straw leaf, straw stem, and whole straw [25], as shown in Table 9.1. When appropriate coupling agents are selected for the fiber treatment, both composite strength and stiffness can be higher than the corresponding values for neat resin. Lei et al. investigated the effect of coupling agent or compatibilizer, i.e. carboxylated PE (CAPE), titanium-derived mixture (TDM), and maleated PE (MAPE) on the mechanical properties of recycled HDPE composites reinforced with 30 wt% pine or bagasse fibers [26]. Results showed that the composites using MAPE were equal or higher in tensile strength, as compared to the recycled neat resin. Storage modulus also increased by a greater extent than the neat resin. Accordingly, the function of a composite is accomplished. On the other hand, impact strength (Izod, notched) decreased compared to the neat resin. Rozman et al. also reported that the strength level of oil-palm-empty-fruit-bunch fiber reinforced PP composites using a coupling agent and compatibilizer was almost equal to the matrix resin, but the impact strength (Izod, unnotched) was lowered [27].

Toughened Green Composites: Improving Impact Properties 221 As mentioned above, many papers clearly point out that the interfacial interaction using coupling agent and/or compatibilizer is necessary for improving static strength and stiffness of the composites. This is especially true for the combination of hydrophilic constituents such as natural fibers and hydrophobic constituents such as olefin polymers. Thus, it can be concluded that, if the minimum necessary interfacial strength is not met, both static and impact strengths of the composites decrease.

9.3.2 A Pattern of Increase in Tensile Strength and Decrease in Impact Strength In this section, green composites, primarily containing natural fibers as green constituent, are discussed from the viewpoint such that fiber length is a significant factor on deciding impact strength. Bos et al. investigated mechanical properties of flax/compatibilized-PP composites produced through injection molding after kneading or extruding process [28]. Both tensile and flexural properties increased with increase in fiber content, but impact strength (Charpy, unnotched) decreased to less than the neat PP resin, as shown in Table 9.2 and Figure 9.3. The average lengths of fibers removed from the resin after injection molding was 0.164 and 0.155 mm for 28 wt% and 51 wt% flax kneaded composites, respectively, and 0.191 and 0.172 mm for 28 wt% and 51 wt% flax extruded composites, respectively. All of these lengths were significantly lower than the critical fiber length of 0.4 mm for the studied fiber/resin combination. They considered that the compatibilizer used in their study induced a strong interface, leading to splitting in the secondary wall of the fiber during the impact test. This phenomenon might be related to low impact properties. Similar trend, showing lower impact properties (Charpy, unnotched) compared to the neat resin, has been reported by Kim et al. as well [29]. Their results are also presented in Table 9.2. Injection molding method was used for long pellets with ramie/PP system of which maleic anhydride-grafted polypropylene (MAPP) was included as a compatibilizer. Results showed that both tensile strength and Young’s modulus of ramie/PP composites increased with increase in fiber content. However, the impact strength was less than that of the neat specimen. This was despite the use of long fiber pellets in the composites. The reason for the low impact strength was explained as the fiber length reduction to 1.56 mm during kneading process from the estimated fiber length of 2 mm in the pellet. However, this length was still longer than the critical fiber length of 0.47 mm [30]. Gironès et al. have reported that flexural strength and modulus of injection-molded abaca/ PP composites (that included MAPP) increased with increase in the fiber

Type of impact test

Charpy (unnotched)

Charpy (unnotched)

Izod (unnotched)

Izod (notched)

Type of fiber

Flax

Ramie

Abaca

Mixture of flax, hemp, and kenaf fibers

Mixer compounding + injection

Extrusion + injection

Extrusion + injection

Extrusion + injection

Production method

77.4 87.3 104

30 40 50 34.2 48.2 61.9

0 5 10 20

Nylon6

Nylon6/NFB

27.8

62.8

55.9

50 20

62.9

40

Abaca/6% MAPP/PP

60.3

30

40.2

51.1

20

0

43.5

10

PP

36.2

0

68

Ramie/PP

58

NMT

67

Kneaded

Extruded

48 51

45

52

NMT

28

Extruded

Kneaded

4.36

3.62

3.25

2.90

6.0

5.2

3.9

2.8

1.1

6.3

5.92

5.41

4.26

3.27

2.35

8.4

6.2

8.5

5.2

5.6

4.6

4.5

5.29

5.79

6.03

7.23

30.3

31.0

32.8

48.4

(impact strength is shown in Figure 9.4)

14.5

18.5

20.3

21.1

23.7

67.9

(impact strength is shown in Figure 9.3)

Note

Ozen, et al. NFB, Natural Fiber Blend, 2013 [32] mixture of flax, hemp, and kenaf fibers Izod impact testing: ASTM D 256

Girones, In the corresponding column et al. flexural strength and modu2011 [31] lus are shown. Elongation at break in this paper is estimated as deflection at break due to the testing method

Kim, et al. (2013 & 2014) [29, 30]

Bos, et al. NMT: Natural fiber Mat 2006 [28] Thermoplastic

Impact Fiber Tensile Tensile Elongation Impact content strength modulus at break strength strength Type of specimen (wt%) (J/m) Ref. (MPa) (GPa) (%) (kJ/m2)

Table 9.2 Relation between tensile and impact properties in green composites.

222 Advanced Green Composites

Charpy (Unnotched)

Flax, Cordenka

Injection molding

Izod Compression (notched and molding unnotched)

Extrusion +compression molding

Flax

Sisal, Charpy banana, jute (unnotched) and flax

50.8 58.0

20 30 40

50.4

54.2

30 10

49.2

20

PLA/Cordenka

44.5 42.7

0 10

35.3

PLA/flax

50

PP/D (90 direction)

40.2

PLA

50

PP/D (0 direction)

28.5

47

45 0

40

30

48

44 43

50

35 20

47

45

40 24

43

33

PP

PP-flax

PP-jute

44

48

43 20

46

34

PP-banana

45

25

PP-sisal

47 49

0 20

PP

1.3

3.11

4.85

3.97

3.27

6.31

5.06

3.90

3.11

5.0

6.5

1.5

4

3

2.3

4.9

4.2

3.2

3.3

3.3

2.3

3.6

3

2.8

2.4

2.2

1.4

24

751*2 450*2

403*1 266*1

51.3

72.2

63.0

43.4

11.1

10.5

9.97

16.1

553*2

24*1

21

27

25

14

20

18

17

15

16

29

27

23

19

Bax, et al. 2008 [35]

Oksman, 2000 [34]

Oksman, et al. 2009 [33]

(Continued)

D: Danflax, *1 notched (J/m), *2 unnotched (J/m) Izod impact testing: ASTM D 256

Toughened Green Composites: Improving Impact Properties 223

Charpy Extrusion + (Unnotched injection and notched)

Charpy (notched)

Jute

Cellulose abaca, jute

30

PP/abaca

35.2

PHBV/Ecoflex/ jute

41.7

30 28.0

27.3

30.2

30 0

23.3

20

PHBV/Ecoflex/ abaca

Yarn coating + PHBV/Ecoflex injection PHBV/Ecoflex/ cellulose

19.5 22.7

PP/EPDM/Jute

15.1

23.0

5

0

PP/EPDM

72

10

0

pure PP

PP/cellulose

29

0

PP 44

92

74

30

PLA/cellulose

63

0

7.0

4.4

4.4

2.1

1.80

1.59

1.56

1.28

0.897

1.10

4.01

4.93

1.50

5.85

8.03

3.37

0.8

0.9

2.3

7.0

0.706

0.909

0.929

0.989

1.25

1.00

4.98

6.18

32.4

4.81

9.54

10.7

22.9

39.4

64.6

30.0

11.1

5.3

3.5

7.9

5.3

2.2

Bledzki, Data plots of impact strength et al. in Fig. 3 of this reference 2010 [40] are converted into numerals by Simple Digitizer. Impact strength at room temperature (23 C) are shown.

5.00* Khalili, EPDM: Ethylene-propyleneet al. diene-monomer 2011 [37] Elongation at break is shown 20.1* as a ratio to pure PP 16.3* Data in the column of impact strength (J/m) are notched 12.1* impact results (Unit: kJ/m2) 5.85* 20.9*

Note

Bledzki, Impact strength data at room et al. temperature (23°C) are 2009 [36] shown.

Impact Fiber Tensile Tensile Elongation Impact content strength modulus at break strength strength 2 Type of specimen (wt%) (J/m) Ref. (MPa) (GPa) (%) (kJ/m )

Yarn coating + PLA injection PLA/abaca

Charpy (notched)

Abaca, cellulose

Production method

Type of impact test

Type of fiber

Table 9.2 Cont.

224 Advanced Green Composites

Toughened Green Composites: Improving Impact Properties 225 content, but the impact strength (Izod, unnnotched) decreased (shown in Table 9.2 and Figure 9.4) [31]. The fiber lengths extracted from the composites gradually decreased with increase in fiber content. The average lengths were 0.723, 0.861, 1.05 and 1.1 mm for 50, 40, 30 and 20 wt% fiber contents, respectively, although these lengths are higher than the critical fiber length of 0.40 mm mentioned above. This result demonstrates that higher fiber contents bring lower fiber lengths and that results in lower impact strength. Furthermore, they found that addition of MAPP did not affect the amount of energy absorbed. A pattern of increase in tensile strength and decrease in impact strength mentioned above is not just limited to the case of PP resin. Ozen et al. [32] reported that tensile strength and Young’s modulus of polyamide (Nylon 6) matrix composites reinforced with mixed natural fibers, i.e. flax, hemp and kenaf, increased with increase in fiber content, but the impact strength (Izod, notched) decreased (see, Table 9.2). A common point in the above four papers is the production method which included the process of compounding followed by injection molding. It is hypothesized that higher fiber content induces more intense twining between fibers during the process, and results in fiber breakage resulting in shortening of the length. On the other hand, Mieck et al. reported that impact strength (notched test, no impact testing method listed) of flax/PP composites increased with increase in fiber length [8]. A significant increase of notched impact strength was realized for l  7–8 mm and

Charpy impact (kJ/m2)

40

30

20

10

0 0

0.1

0.2 0.3 0.4 Fibre weight fraction (–)

0.5

0.6

Figure 9.3 Unnotched Charpy impact of kneaded flax/PP/MAPP and kneaded flax/ PP, extruded flax/PP/MAPP and extruded flax/PP compounds, compared with hackled flax/PP/MAPP NMTs, versus fiber weight fraction. The fitted lines are: — kneaded and --- extruded flax/PP/MAPP. The PP value is for Retiflex PP used for the NMTs, the PP used for the compounded materials did not break during the test. Error bars falling within the markers are not shown [28].

226 Advanced Green Composites higher fiber content. For a non-woven composite with an accomplishable fiber length l = 30 mm and fiber content more than 40 wt%, as shown in Figure 9.5, impact strength is higher than 20kJ/m2. Oksman et al. also confirmed that a pattern of increase in tensile strength and decrease in impact strength was seen in PP resin composites reinforced with sisal, banana, jute and flax fibers [33]. It should be noted that in their study MAPP was not used. The typical fiber length value extracted from the composites for sisal fibers was 3–5.5  mm and for banana fibers was 1.5–2  mm. The flax and jute fibers were shorter and the typical value was 0.75 mm, which is also larger than lc mentioned above. On the other hand, Oksman [34] reported in an earlier study that natural fibers such as flax, improved the impact

Impact strength (kJ/m2)

12 10 8 6 4 20 % abacá 30 % abacá 40 % abacá 50 % abacá

2 0 0

2

8

4 6 % MAPP

10

Figure 9.4 Impact strength of abaca-reinforced composites with and without MAPP coupling agent [31].

Notched impact strength (kJ/m2)

40

30

20

Fibre diameter: 60 m Fibre modulus: 48 kN/mm2 Fibre tenacity: 660 N/mm2 22 wt% 31.9 wt% 40.2 wt% 50.6 wt%

10

0 0.01

0.1

1

10

100

Fibre length l (mm)

Figure 9.5 Notched impact strength as a function of fiber length in flax/PP composites [8].

Toughened Green Composites: Improving Impact Properties 227 behavior of PP, especially when long fibers (supercritical fiber) were used. In this report, impact strengths of notched and unnotched specimens using needle punched nonwoven flax/PP increased higher than impact strength of the neat resin (results of Danflax version only is reported in Table 9.2). Thus, change in impact strength of green composites with change in fiber content is closely related to the degree of fiber length. It can therefore, be concluded that the saturating fiber length for impact strength, mentioned earlier, is longer than the critical fiber length, similar to glass fiber-reinforced composites. According to Bax and Müssig [35], on the contrary, impact strength (Charpy, unnotched) of flax fiber/PLA composites increased with increase in fiber content, although the increasing ratio was not so remarkable, as compared to those of tensile strength and modulus, as shown in Table 9.2. However, they did not explore the effect of fiber length on the impact strength, which they listed as a future research subject. In addition, they made a comment that the magnitude of the impact strength is less than the neat resin. On the other hand, Bledzki et al. [36] reported that impact strengths (Charpy, notched) of abaca/PP and abaca/PLA became higher than the neat resins, respectively, at both –30 C and room temperature. This difference was clearly dependent on whether the specimens possessed the notch or not. Based on these results, it can be concluded that notched natural fiberreinforced composites show more toughening effect than notched resins.

9.3.3 Effect of Toughened Resin In natural fiber reinforced composites, there are also some studies where natural or synthetic elastomer particles are added to the resin to improve its toughness, as mentioned earlier in section 1 for glass fiber reinforced composites. According to Khalili et al. [37], addition of ethylene-propylenediene-monomer (EPDM) to PP is very effective for improving its impact properties. Although, tensile strength decreased when EPDM was added, Charpy impact strength of unnotched specimens increased 2.15 times, and that of notched specimens increased 4.17 times higher than the neat resin without EPDM. On the other hand, both impact strengths decreased when jute fibers were added to PP/EPDM. Finally, impact strength of 30% jute fiber reinforced composites decreased by 15% and 28% in unnotched and notched specimens, respectively, compared to PP/EPDM. When impact strengths of 20% and 30% wt% jute fiber reinforced composites fabricated by Khalili et al. [37] are compared with 24% wt% jute fiber-reinforced composites made by Oksman et al. [33], shown in Table 9.2, it is clear that addition of EPDM to PP is not so effective. Not only the tensile strength and modulus

228 Advanced Green Composites values are lower, but also the impact strength is lower. Similar effect of EPDM addition was also seen in vetiver grass fiber/PP composites [38]. They added EPDM or natural rubber to PP in order to improve impact strength of natural fiber based composites. Results showed that the impact strength (Izod, unnotched) increased with increase in rubber content, but it is less than neat PP resin, even when 30 wt% EPDM or natural rubber was used. The addition of plasticizer has also been tried for PLA based green composites to improve their toughness. According to Oksman et al., triacetin (glycerol triacetate ester), which usually improves the elongation of PLA from less than 10% to more than 250%, did not show any positive effect on the impact strength (Charpy, unnotched) of flax fiber-reinforced PLA composites [39]. Bledzki et al. reported that the impact strengths of jute and abaca fiber composites did not improve significantly, as shown in Table 9.2 [40]. They used more ductile resin, a blend of poly (3-hydroxybutyrate-co3-hydroxyvalerate) (PHBV) and poly (butylene adipate-co-butylene terephthalate) (Ecoflex), as the resin, but the impact strength (Charpy, notched) was almost equal to the levels of PLA and PP based composites. Thus, it may be concluded from the results obtained up-to-date that toughened resins do not necessarily improve the impact strength of natural fiber composites. As a result, some innovative studies such as -modification developed in the field of glass fiber composites can be expected to be helpful.

9.3.4 Approaches to Increase Both TS and IS Although Oksman, et al. reported decrease in impact strength of sisal/PP, banana/PP, jute/PP and flax/PP composites, as mentioned above, they left a clue that toughness improvement such as the degree of impact strength also depended on the fiber type [33]. Their results indicated that the sisal/ PP composites had the best impact properties although the level was only slightly higher than the neat PP resin. These composites, furthermore, maintained the longest fibers after extrusion. They made a comment that this depends on fiber toughness (sisal has best elongation to break) and long fiber pullouts, which implies that tough fibers can be maintained in an extended state in the matrix during the composite production process, and this can lead to effective fiber pullout. Such a different tendency between tensile and impact strength becomes more clear through the study by Bledzki, et al. [36]. They reported mechanical properties of PLA matrix composites with abaca and regenerated cellulose fibers processed by using combined molding technology, in which filament yarn is coated with the resin out of the extruder. Through this technology, most of the fibers remain longer than 2 mm after compounding.

Toughened Green Composites: Improving Impact Properties 229 Although being shortened below 2 mm after injection molding, the fibers maintained a relatively long state. In this state, the fiber diameter is also reduced and, therefore, high aspect ratio is still maintained. As a result, impact strength (Charpy, notched) of abaca fiber/PLA composites was improved many folds to 5.3kJ/m2 in average from 2.2 kJ/m2 obtained for the neat PLA at room temperature, as shown in Table 9.2. Furthermore, regenerated cellulose fiber composites showed the average impact strength of 7.9kJ/m2. They estimated that this significant improvement was attributed to more pull-outs in regenerated cellulose fiber composites because of the smooth surface of the cellulose fibers, different from irregular crosssections of abaca fibers. They added that such an improvement in impact properties was further confirmed when PP was used as the resin. Bledzki, et al. reported furthermore that, when more ductile resin, a blend of PHBV and Ecoflex, was used, the impact strength (Charpy, notched) of regenerated cellulose fiber composites was dramatically improved, although this resin gave lower tensile strength than PLA and PP composites, and was not able to improve impact strength of natural fiber composites, as stated earlier [40]. Based on the results discussed above it can be concluded that regenerated cellulose fibers result in improvement of impact strength which may be a result of their smooth surface that facilitates fiber pull-out.

9.4 Role of Large Elongation at Break in Regenerated Cellulose Fibers As mentioned above, regenerated cellulose fiber reinforced PLA composites exhibited good impact properties. Since the cellulose fibers have a good elongation at break, this property may also be related to improvement of impact properties. Cordenka 700 single fibers used in the above paper by Fink and Ganster showed an average of 13% elongation at break as well as tensile strength and Young’s modulus of 833 MPa and 20 GPa, respectively [41]. According to their paper, Enka viscose also showed high elongation at break of 24%, although its tensile strength is only 308 MPa, less than half of that of Cordenka 700. Such a high elongation at break together with high strength has never been achieved in plant-based natural fibers. However, these strength and modulus values are comparable to some natural fibers such as flax. Bax and Müssig also reported that, impact strength (Charpy, unnoched) of Cordenka/PLA composites increased by 3.45 to 6.5 times those of flax/PLA composites, although tensile strength of the former was only slightly higher than the latter, i.e. 1.03 to 1.18 times higher (Table 9.2) [35]. The highest impact strength achieved 72.2 kJ/m2 for

230 Advanced Green Composites 30 wt% Cordenka/PLA composites, an excellent value, as compared to 16.1 kJ/m2 of the neat PLA resin. According to Ganster et al., mechanical property level of PP based composites using Cordenka 700 yarns is comparable with short glass fiber reinforced polypropylene (GF/PP) composites with advantages in tensile strength as well as impact strength [42]. The stiffness of the composites is lower than for GF/PP due to the anisotropic nature of the fiber, as shown in Table 9.3. They added that other thermoplastics like PE, high impact polystyrene and PLA have been successfully reinforced as well. Thus, it could be said that Cordenka is a great reinforcing material that also results in good toughness. Bax and Müssig have commented that (i) debonding, (ii) pull-out and (iii) fiber fracture are three mechanisms of energy absorption during impact [35]. Below, we consider which energy is the most influential for the absorption energy and discuss the reasons why Cordenka fiber-reinforced composites bring the highest impact strength. Matthews and Rawlings pointed out that critical strain release energy of fibers was four times larger than that of debonding at the fiber/matrix interface [7]. Furthermore, the energy spent during pull-out is much more than that of debonding. Their calculations showed that the most influential energy absorption is caused by pull-out mechanism (See appendix). However, these ratios were calculated for a fiber along the loading axis. In many actual composites, however, the fibers are randomly aligned two- or even three-dimensionally. According to Li et al. [43] and Wu et al. [44], the inclined discontinuous fibers embedded in the matrix at an angle from a crack surface sustain a higher stress than the fiber perpendicular to the crack plane ( ) (see, Figure 9.6). This is called the “snubbing effect”, and verified experimentally through fiber-reinforced cement composite. This concept was also extended to glass fiber reinforced polymer composites Table 9.3 Comparison of selected mechanical characteristics of glass fiber reinforced PP (GF-PP, Piolen G30CA60) and Cordenka-PP composite with 30 wt% fibers (PPRayCo30) [42]. Property

Unit

Value for Value for GF-PP PPRayCo30

Tensile strength

MPa

56.3 ± 0.4

78.7 ± 1.1

Tensile modulus

GPa

Charpy unnotched impact strength, 23 °C

kJ/m

4.1 ± 0.2

2.94 ± 0.04

2

35± 3

89 ± 8

2

39 ± 3

88 ± 6

2

13.4 ± 0.3

12.1 ± 0.7

2

8.5 ± 0.3

9.3 ± 1.5

Charpy unnotched impact strength, 18 °C kJ/m Charpy notched impact strength, 23 °C Charpy notched impact strength, 18 °C

kJ/m

kJ/m

Toughened Green Composites: Improving Impact Properties 231 Matrix

Fibre

P

s l

Figure 9.6 Pull-out of an oblique fiber from a matrix (l: embedded length, s: sliding length) [45].

by Fu et al. [45]. Relation between two loads P and P , respectively, occurring in the fibers embedded in the matrix and out of the crack surface is given as P  = Pexp(f ). Where, P = P =0 and f is the snubbing coefficient, which for nylon and PP fibers were obtained as 0.994 and 0.702, respectively. Since the term exp(f ) is higher or equal to 1, P must be higher than P =0. In the combination of lower elongation-at-break fibers such as natural fibers and polymeric resins, this theory cannot be applied necessarily, but it may be applicable at relatively low angles. In a large angle range, the fibers would be broken in bending mode [46]. In any case, the fiber load in the matrix is less than P , when a crack making an angle with the fiber axis passes through. From such mechanism, even if the embedded length l is slightly less than lc/2 (lc: critical fiber length), the fiber would be possible to fracture at the crack surface more easily rather than the occurrence of pullout. In other words, fiber fracture mechanism becomes more dominant, as compared to the situation of fibers embedded perpendicular to crack. In this situation, in order to toughen the composites, the fibers must have high toughness. High toughness of regenerated cellulose fibers such as Cordenka, as mentioned above, can therefore result in high impact strength, as compared to natural fiber composites.

9.5 Toughened Cellulose Fibers and Green Composites 9.5.1 Toughening Mechanism of Regenerated Cellulose Fibers Next question is to ask why does regenerated cellulose fibers exhibit a higher toughness, and plant-based natural fibers show a low elongation-at-break?

232 Advanced Green Composites 0.8 CA

CB

Stress (GPa)

0.6 CD 0.4

CC

0.2

0.0

0

2

4

6

12 8 10 Strain (%)

14

16

18

Figure 9.7 Stress-strain curves for Cordenka EHM (CA), Cordenka 1840 (CB), Enka viscose (CC), and Lyocell (CD) regenerated fibers [47].

Eichhorn et al. reported on the deformation processes of four kinds of regenerated cellulose fibers [47]. As is known, the crystalline structure of regenerated cellulose fibers is Cellulose II, which results in 90 GPa of crystalline elastic modulus, while Cellulose I, the main constituent of plantbased natural fibers, exhibits 140 GPa, approximately 1.5 times larger than Cellulose II. Thus, Young’s modulus of regenerated cellulose fibers is lower than that of plant-based natural fibers. Eichhorn, et al. first performed tensile tests of the fibers. Typical stress-strain curves of regenerated cellulose fibers are shown in Figure 9.7, in which most fibers show larger elongations at break than plant-based natural fibers [47]. The present author calculated the areas under stress-strain curves, i.e. the work of fracture, by converting the plots in Figure 9.7 into numerals by Simple Digitizer, and pasting them to Microsoft Excel 2010. Results show that these areas occupied respectively 16.1, 46.3, 25.6 and 24.7 MPa for Cordenka EHM (CA), Cordenka 1840 (CB), Enka viscose (CC), and Lyocell (CD), respectively. On the other hand, the work of fracture of glass and flax fibers are estimated as 23.1 and 10 MPa per unit volume, by assuming for simplicity, that their deformation behaviors are linear elastic, and it is given as 2 / 2E. (Tensile strengths of glass and flax fibers are given as 1800 MPa and 1000 MPa, and Young’s moduli are given as 70 and 50 GPa, respectively.) Thus, it is clear that the work of fracture in CB (Cordenka 1840 type fibers) is much higher than the other fibers. According to the same paper by Eichhorn et al., CB showed average of 12.7% of elongation at break, tensile strength of 660 MPa and Young’s modulus of 16.9 GPa [47]. Therefore, Cordenka 700 mentioned in the previous section is estimated to have higher work of

Toughened Green Composites: Improving Impact Properties 233 Fibrils

Fibrils Fiber axis

A C A C A C A C A C A C A C A C A C A C A C A C A C A C A C A C A C A C A C A C

Amorphous domains

Crystallite

A C A C A C A C A C A C A C A (a)

(b)

Figure 9.8 Schematic of (a) modified series model, (b) possible physical structure of a semicrystalline cellulose fiber [47].

fracture (toughness). Next, Eichhorn et al. investigated the shift in Raman spectra wavenumber during tensile test [47]. They found that upon straining, the peak position at 1095 and 895 cm-1 bands shifted toward lower wavenumbers, and the degree of the shift depended on the level of strain hardening of the fibers. Slope of the shift width to strain is called Raman band shift sensitivity, which depends on the Young’s modulus (stiffness) of the fibers. They considered from a modified series-aggregate model by Yeh and Young [48] that fibers are built up of parallel arrays of identical fibrils, which contain crystalline domains interrupted by amorphous domains, as shown in Figure 9.8. A modified series aggregate model has been developed to explain the non-linear elastic response of highly oriented fibers. This model gives a full explanation of the stress–strain behavior of aramid, PET and cellulose fibers. However, their discussion that should be noted is the behavior of the fibers after yielding, rather than non-linear elastic behavior [47]. Once the fiber yields, slippage of the crystalline domains in the fibrils, past each other, is thought to dominate. Therefore, if the fiber includes more crystalline parts, the stress-transfer becomes better between neighboring portions of the fibrils. This primary mechanism develops high strain hardening in fibers. Physically, it represents a high deformation resistance during plastic deformation. It can be concluded, finally, that Cordenka fibers behave as a tough bridging medium between crack surfaces during the impact tests, resulting in high impact strength.

234 Advanced Green Composites

Crystalline cellulose

Amorphous cellulose, mannose, galactose, xylose, arabinose, etc.

Crystalline cellulose Hemicellulose

Figure 9.9 Chemical structure of untreated ramie fiber microfibrils [50, 52].

9.5.2 Mercerization Effect Can toughness increase mechanism be brought to plant-based natural fibers? Mercerization is a method of alkali-treatments for plant-based natural fibers. Mercerized fibers have a Cellulose II crystal structure, but still maintain the form of an original natural fiber. According to Bledzki, et al., mercerized hemp yarn exhibits much higher elongation at break of 6.86% [49]. Goda, et al. also reported that mercerized ramie single fibers showed a large fracture strain of 7.2% when a load during mercerization was applied [50]. Such a large deformation mechanism is described below. Natural fibers can be considered as a bundle of microfibrils bound together by amorphous lignin, waxy materials and other impurities. Microfibrils are made up of long cellulose chains and hemicellulose [51]. Figure 9.9 shows schematic representation of untreated ramie fiber monofilament showing the structure of microfibril bundle [50, 52]. The microfibrils are extensively hydrogen bonded with the crystalline form of cellulose. The detailed chemical structure of the untreated microfibril is shown in the right-hand side of Figure 9.9 [52]. It contains long crystalline portions composed of cellulose as well as small amorphous regions composed of several pentose and hexose sugars. Upon application of tensile stress, the microfibrils in the untreated fiber tend to slip past one another. However, the hemicellulose termed as ‘Biocement’ present in between the microfibrils tends to retain the microfibrils in their original positions. This results in the reversible behavior of microfibrils, which is responsible for the less plastic deformation in untreated fiber. When the applied stress on the fiber increases, a small scale

Toughened Green Composites: Improving Impact Properties 235

Untreated cellulose microfibrils

(a) Mercerization without load application

(b) Mercerization with load application

Hemicellulose

Figure 9.10 Schematics of cellulose microfibril structure before and after mercerization (a) without load application and (b) with load application [50].

of irreversible slippage may occur within the microfibril. However, presence of hemicellulose between microfibrils acts as a bonding agent and restricts the slippage to an extent. Hence, the untreated fiber behaves like an elastic body up to fracture, though it is not elasticity in the strictest term. The change in the deformation behavior of the treated fibers is attributed to various changes that occur in the physical and chemical structure of the fibers occurring during mercerization. Cellulose molecular chains in a microfibril lose their crystalline structure locally by mercerization. At the same time the alignment of the microfibrils is destroyed and the overall crystallinity reduced due to the breakage of the three dimensional network of cellulose by the extensive cleavage of hydrogen bonding, as shown in Figure  9.10. The decreased crystallinity and the disorder of the microfibrils alignment decrease the stiffness of the fiber. Mercerization also removes hemicellulose from the fiber resulting in an easy deformation of the cellular networks. The extensive hydrogen bonding network will also be broken and the ordered structural arrangement of cellulose may loosen up. The deformation of individual microfibrils becomes easier due to the absence of interlocking hydrogen bonding. Through the above-mentioned changes, the microcrystalline microfibrils may be more prone to irreversible slippage and even

236 Advanced Green Composites breakage resulting in the plastic deformation of the fiber upon tensile load. When the stress acts upon the slipped out microfibril, large-scale irreversible slippage may occur. Thus, higher elongation at break is observed for the mercerized fibers because of the decreased crystallinity in cellulose and the loosely bound structure of the microfibrils within the fiber. Effects of mercerization on the impact properties of ramie-fiber-reinforced green composites were explored through the drop-weight impact testing method (ASTM D3763) by Suizu et al. [53]. Table 9.4 shows maximum loads and impact energies of the laminated composites with unidirectional and plain-woven fabric layers using untreated and mercerized ramie yarns. These laminated composites are respectively denoted as untreated (UT) laminate and alkali-treated (AT0) laminate. As shown in Table 9.4, the projectile, with potential energy of 7.36 J, penetrated the UT laminate and the impact energy was 5.71 J. However, these AT0 laminate specimens were not penetrated by the projectile with same potential energy. The projectile potential energy was then further increased to 24.5 J, and the impact test was carried out using a new specimen of AT0 laminate. In this case, AT0 laminate was penetrated with the impact energy of 12.7 J. This clearly indicates that improvement through mercerization of ramie yarns resulted in increasing their impact properties to almost twice as high as those of laminated composites reinforced with UT yarns. AT0 fabric laminates also exhibited better impact properties, as shown in Table 9.4. For comparison, glass fiber roving cross mat reinforced unsaturated polyester laminates (GFRP) were fabricated by hand lay-up method. These laminates were also impact-tested in the same way. The obtained absorbed energy is also shown in Table 9.4. It can be seen that AT0 fabric laminate is higher in absorbing energy than GFRP, though the thicknesses of both materials were different. Figures 9.11 (a), (b) and (c) respectively show the representative damage morphology of UT and AT0 laminates, and AT0 fabric that suffered

5 mm

10 mm (a)

(b)

5 mm (c)

Figure 9.11 Impact damage morphology on the reverse side of green composites reinforced by (a) UT yarns, (b) AT0 yarns and (c) AT0 fabric [53].

1.34

62.0

45.0

57.1

55.9

42.7

Fiber volume fraction (%)

24.5

24.5

1.95

1.90

0.699

penetration

penetration

penetration

rebound

penetration

Maximum load Impact (kN) response

*1 UT: untreated, *2 AT0: alkali-treated (mercerized without load application), *3 Glass fiber roving cross mat

2

GFRP*3

1.89

2.21

1

2

2.33

2.24

Thickness (mm)

3

2

AT0 Fabric*2

AT0 laminate*2

UT laminate

*1

Number of samples

Table 9.4 Impact properties of laminated composites reinforced with mercerized ramie yarns [53].

9.04

13.9

12.7

–

5.71

Impact energy (J)

Toughened Green Composites: Improving Impact Properties 237

238 Advanced Green Composites penetration [53]. The UT laminate is cruciformly damaged with yarn breaks perpendicular to the yarn direction, as shown in an arrow of Figure 9.11 (a), which means that the yarn cannot be deformed greatly. In contrast, the AT0 laminate is damaged with an interfacial crack of about 60  mm length between yarn fibers without any yarn breakage, which signifies that the yarn absorbs more impact energy because of toughness improvement for ramie yarns. The AT0 fabric also exhibited a higher impact property because of its plain-woven structure as well as mercerization effect. Consequently, the best merit of mercerization was seen in the form of significant improvement in impact properties when mercerized natural fibers were used as reinforcement of green composites.

9.5.3 Other Beneficial Chemical Treatments Beneficial chemical treatments for impact strength improvement, other than mercerization, have been studied as well. Kabir, et al. reported that chemical treatment by benzenediazonium salt in alkaline medium improved impact strength of flax/PP composites [54]. They treated jute fibers with o-hydroxybenzenediazonium salt (o-HBDS) separately, by adding 0.02 moles (3.04 g) of o-amino phenol, NaNO2, and HCl. Through this treatment, impact strength of o-HBDS-treated in alkali medium jute fiber-PP composites were approximately two times higher than that of pure PP. These results are shown in Table 9.5 and indicate that the treatment of jute fibers with o-HBDS in alkali medium increased the compatibility between jute fiber and PP. This may be due to coupling of o-HBDS with hydroxyl groups in jute fibers, thereby reducing their hydrophilic nature. The treatment of jute fibers with o-HBDS may also increase their interfacial bonding with PP. They estimated that a stronger force was required to pull fibers out from the composites. However, they did not mention the fiber length used. As discussed earlier, when the interfacial bond increases, the critical fiber length decreases. As a result, as known from the discussion mentioned above, fiber pull-out is hard to occur, and fiber fracture starts to occur. From these results, it may be concluded that toughening of jute fibers occurs through o-HBDS treatment. Xia, et al. studied the effect of fiber surface modification by alkaline solution, maleic anhydride (MAH) and silane (KH550) on the impact strength of flax/PLA composites [55]. PLA/alkali-silane KH550-treated flax fiber reinforced composites showed the highest impact strength, which was 20% higher than the neat PLA. They estimated that the higher impact strength was a result of strong interface formed by the chemical treatment. On the other hand, they discussed another factor for the toughening effect of flax/PLA composites in terms of transcrystals being formed around the fibers. They concluded that

Type of Type of Production fiber impact test method Type of specimen Jute Charpy Extrusion + PP (notched) injection Raw jute/PP alkali-treated jute/PP o-HBDS-treated/PP Raw jute/PP alkali-treated jute/PP o-HBDS-treated/PP Raw jute/PP alkali-treated jute/PP o-HBDS-treated/PP Raw jute/PP alkali-treated jute/PP o-HBDS-treated/PP Flax Izod Extrusion + PLA (unnotched) injection PLA/UF PLA/AF PLA/MF PLA/KF

Fiber Tensile Tensile content strength modulus (wt %) MPa) (GPa) 0 28.3 1.06 20 25.9 1.78 30.1 1.99 45.3 3.54 25 25.8 1.76 29.4 1.96 44.7 3.45 30 25.6 1.76 28.6 1.93 43.8 3.35 35 25.1 1.75 28.1 1.91 43.0 3.34 5 54.1 2.68 59.5 3.38 58.4 3.26 57.6 2.78 58.1 3.00

Elongation Impact at break strength (%) (kJ/m2) Ref. Note 15.0 3.50 Kabir, Data plots in Figs. 1, 2 et al. and 5 of this refer1.36 2.51 2010 ence are converted 1.51 2.90 [54] into numerals by 1.48 6.49 Simple Digitizer. 1.33 2.44 1.50 2.79 1.46 6.11 1.28 2.33 1.49 2.72 1.44 5.54 1.22 2.10 1.48 2.61 1.41 5.08 13.3 Xia, UF: PLA/untreated et al. fiber composite 14.8 2015 AF: PLA/alkali-treated 15.1 [55] fiber composite 14.4 MF: PLA/alkali-maleic 16.1 anhydride-treated fiber composite KF: PLA/ alkali-silane KH550-treated fiber composite

Table 9.5 Mechanical properties of green composites reinforced with natural fibers undergone various chemical treatments.

Toughened Green Composites: Improving Impact Properties 239

240 Advanced Green Composites crazing and/or crack propagation were being hindered by fibers, energy being absorbed through fiber fracture and transcrystallization were the main reasons for toughening along with good compatibility between fibers and matrix and uniform distribution of fibers in the matrix. On the other hand, they observed pulling and tearing states of impact fracture surfaces. Such fracture patterns of microfibrils in the fibers are also considered to be related to improvement in toughness. And this degree of pulling and tearing mechanism may also be related to the kinds of chemical surface treatments given to fibers. Thus, while green composites have been toughened through several chemical treatments of natural fibers, exact toughening mechanisms have not been discussed sufficiently yet. As a result, further studies on toughening green composites are needed that can shed more light on the toughening mechanisms so these composites may compete with conventional composites in practical applications.

9.6 Conclusions In this chapter, the author has focused on impact strength of green composites. The main conclusions are summarized as follows: 1. If the minimum necessary natural fiber/resin interfacial strength is not met, both static and impact strengths decrease. 2. Optimal fiber length necessary to obtain good impact strength in green composites is more than the critical fiber length, similar to the glass fiber-reinforced composites. 3. It can be concluded from the results obtained up-to-date that toughened resins cannot improve the impact strength of natural fiber composites. Therefore, some innovative studies such as -modification, developed in the field of glass fiber composites, are needed. 4. High toughness of regenerated cellulose fibers such as Cordenka is concluded to bring a high impact strength, as compared to green composites reinforced natural fibers. It is hypothesized that Cordenka fibers behave as a tough bridging medium between crack surfaces during the impact test. 5. When mercerized natural fibers are used as reinforcement in green composites, the fibers contribute to a significant improvement in composite impact properties. 6. Green composites have been toughened through several chemical treatments of natural fibers as well. However,

Toughened Green Composites: Improving Impact Properties 241 further studies on toughening green composites are needed that can shed more light on the toughening mechanisms. In the toughness improvement area, glass fiber-reinforced polymer composites have made large advances from both experimental and theoretical points of view. The area of green composites should find many different toughening mechanisms as well as learning the points, as mentioned in this chapter. It should be noted, however, that the degree of toughness cannot be evaluated only from impact properties, but should be discussed from many other evaluation parameters, such as interlaminar fracture toughness [56] and energy absorption [57]. Since these subjects are very much related to evaluation of primary structural materials, such as carbon fiber-reinforced polymer composites, these do not become major research target for green composites, which are yet to be used for primary structures. However, as mentioned in another chapter, advanced green composites can be strong and stiff and hence, can be used in primary structural applications and many more toughness related studies can be expected in the near future.

Appendix The energy of debonding per fiber WD can be calculated by equating it to the strain energy released in the fiber as the stress relaxes [7]. The debonding takes place over a length l of a fiber of diameter df, and therefore WD is given, by assuming a circular cross-section of the fiber, as follows:

WD

W (volume of fiber debonded )

where W is the strain energy in the fiber per unit volume and expressed by the form 2/(2Ef), where, is the fiber stress and Ef is the Young’s modulus of the fiber. By taking into account the volume of the debonded section of the fiber, the energy of debonding per fiber can be obtained as follows [5]:

WD

D2 8E f

l

0

4 l d

2

dl

(9.1)

where, D is a fiber diameter, l is the distance from the matrix surface to the debonding end and is the interfacial shear stress (or strength). l can be estimated as D /4 from the shear-lag model on the condition that interfacial shear stress is given as a constant. When reaches the breaking strength

242 Advanced Green Composites of the fiber f, the maximum energy of debonding occurs. Substituting l = D f/4 for eq. (9.1) and integrating we have

D2

WD ,max

2 f

D2

l

24 E f

2 f

lc

(9.2)

48 E f

where, lc is the critical fiber length. Pull-out energy, Wp, has been calculated by integrating (shear force corresponding to fiber-resin contact area) times (differential sliding distance), as follows [6]:

Dl 2

l

WP

D(l z ) f dz

f

(9.3)

2

0

where, f is the shear frictional stress during pull-out or Coulomb’s frictional stress caused by a force in the radial direction. When l = lc /2, eq. (9.3) gives the maximum pull-out energy, WP, max, as follows:

WP ,max

Dlc 2

f

(9.4)

8

The mean, Wp, mean, of a population with length zero to lc /2 can also be calculated by:

WP ,mean

1 lc / 2

lc /2

0

Dl 2 2

f

dl

Dlc 2 24

f

(9.5)

This is because the location of crack plane varies from the center of the fiber to the end, even if the fiber length is lc. Thus, when f is assumed almost equal to , the ratio of Wp, max to WD, max is:

WP ,max WD ,max

3E f

(9.6)

Assuming that these energies are expended around the fiber breaking stress, can be replaced for the fiber strength f*. When fiber strengths of glass, carbon and flax are roughly given as 1800 MPa, 3000 MPa and 1000 MPaand Young’s moduli of these fibers are given as 70 GPa, 250 GPa

Toughened Green Composites: Improving Impact Properties 243 and 50 GPa, respectively, and the ratios Ef/ f* are roughly in the range of 40 – 80. Therefore, eq. (9.6) can be written as follows:

WP WD

(120 240)

(9.7)

From such estimation, one can understand that the energy of pull-out is much more than the energy of debonding. That is to say, the former energy is a significant parameter to toughen fibrous composite materials, when the fiber length is less than lc. One should understand that, on the other hand, the toughening mechanism depends largely on the magnitude of frictional shear stress, f, at the interface, as shown in eqs. (9.4) and (9.5). When the fiber length is more than lc, by considering a probability that the crack plane crosses uniformly every part along the fiber-axis, the pullout energy is given as:.

WP ,mean

Dlc 2 24

f

lc l

(9.8)

From eq. (9.8) we understand that the pullout energy decreases in proportion with 1/l with increase in the fiber length.

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Toughened Green Composites: Improving Impact Properties 245 24. J. Z. Lu, Q. Wu, I. I. Negulescu and Y. Chen, The Influences of Fiber Feature and Polymer Melt Index on Mechanical Properties of Sugarcane Fiber/ Polymer. Compos. J. Appl. Polym. Sci. 102, 5607–5619, 2006. 25. F. Yao, Q. Wu, Y. Lei and Y. Xu, Rice straw fiber-reinforced high-density polyethylene composite: Effect of fiber type and loading, Industrial Crops and Products, 28, 63–72, 2008. 26. Y. Lei, Q. Wu, F. Yao and Y. Xu, Preparation and properties of recycled HDPE/ natural fiber composites, Composites: Part A, 38, 1664–1674, 2007. 27. H. D. Rozman, G. B. Peng and Z. A. M. Ishak, The effect of compounding techniques on the mechanical properties of oil palm empty fruit bunch– polypropylene composites, J. Appl. Polym. Sci. 70, 2647–2655, 1998. 28. H. L. Bos, J. Mussig and M. J. A. Oever, Mechanical properties of short-flaxfibre reinforced compounds, Compos: Part A 37, 1591–1604, 2006. 29. H. -B. Kim, K. Goda and K. Aoki, Impact properties of short ramie/PP composite materials, in: Proceedings of the 8th International Conference on Green Composites, pp. 8–9, 2014. 30. H. -B. Kim, K. Goda, J. Noda and K. Aoki, Developing Simple Production of Continuous Ramie Single Yarn Reinforced Composite Strands, Adv. Mech. Eng. (Open Journal) 2013, Article ID 496274, 7 pages, 2013. 31. J. Gironès, J.P. Lopez, F. Vilaseca, J. Bayer R., P.J. Herrera-Franco and P. Mutjé, Biocomposites from Musa textilis and polypropylene: Evaluation of flexural properties and impact strength, Compos. Sci. Technol. 71, 122–128, 2011. 32. E. Ozen, A. Kiziltas, E.E. Kiziltas and D.J. Gardner, Natural Fiber Blend— Nylon 6 Composites, Polym. Compos. 34, 544–553, 2013. 33. K. Oksman, A. P. Mathew, R. Långström, B. Nyström and K. Joseph, The influence of fibre microstructure on fibre breakage and mechanical properties of natural fibre reinforced polypropylene, Compos. Sci. Technol. 69, 1847–1853, 2009. 34. K. Oksman, Mechanical properties of natural fibre mat reinforced thermoplastic, Appl. Compos. Mater. 7, 403–414, 2000. 35. B. Bax and Jörg Müssig, Impact and tensile properties of PLA/Cordenka and PLA/flax composites, Compos. Sci. Technol. 68, 1601–1607, 2008. 36. A. K. Bledzki, A. Jaszkiewicz and D. Scherzer, Mechanical properties of PLA composites with man-made cellulose and abaca fibres, Compos: Part A 40, 404–412, 2009. 37. S. M. R. Khalili, R. E. Farsani and S. Rafiezadeh, An experimental study on the behavior of PP/EPDM/JUTE composites in impact, tensile and bending loadings, J. Reinf. Plas. Compos. 30, 1341–1347, 2011. 38. Y. Ruksakulpiwat, J. Sridee, N. Suppakarn and W. Sutapun, Improvement of impact property of natural fiber–polypropylene composite by using natural rubber and EPDM rubber, Compos: Part B 40, 619–622, 2009. 39. K. Oksman, M. Skrifvars and J. -F. Selin, Natural fibres as reinforcement in polylactic acid (PLA) composites, Compos. Sci. Technol. 63, 1317–1324, 2003. 40. A. K. Bledzki and A. Jaszkiewicz, Mechanical performance of biocomposites based on PLA and PHBV reinforced with natural fibres – A comparative study to PP, Compos. Sci. Technol. 70, 1687–1696, 2010.

246 Advanced Green Composites 41. H. -P. Fink and J. Ganster, Novel Thermoplastic Composites from Commodity Polymers and Man-Made Cellulose Fibers, Macromol. Symp. 244, 107–118, 2006. 42. J. Ganster, H. -P. Fink and M. Pinnow, High-tenacity man-made cellulose fibre reinforced thermoplastics – Injection moulding compounds with polypropylene and alternative matrices, Compos: Part A 37, 1796–1804, 2006. 43. V. C. Li, Y. Wang and S. Backer, Effect of inclining angle, bundling and surface treatment on synthetic fibre pull-out from a cement matrix, Compos 21, 132–140, 1990. 44. H. -C. Wu and V. C. Li, Snubbing and Bundling Effects on Multiple Crack Spacing of Discontinuous Random Fiber-Reinforced Brittle Matrix Composites, J. American Ceramic Society 75, 3487–3489, 1992. 45. S. -Y. Fu and B. Lauke, The fibre pull-out energy of misaligned short fibre composites, J. Mater. Sci. 32, 1985–1993, 1997. 46. J. Morton and G. W. Groves, The cracking of composites consisting of discontinuous ductile fibres in a brittle matrix- effect of fibre orientation, J. Mater. Sci. 9, 1436–1445, 1974. 47. S. J. Eichhorn, R. J. Young and W. -Y. Yeh, Deformation processes in regenerated cellulose fibers, Textile Res. J. 71, 121–129, 2001. 48. W. -Y. Yeh and R. J. Young, Molecular deformation processes in aromatic high modulus polymer fibres, Polymer 40, 857–870, 1999. 49. A. K. Bledzki, H. -P. Fink and K. Specht, Unidirectional hemp and flax EPand PP-Composites: Influence of defined fiber treatments, J. Appl. Polym. Sci. 93, 2150–2156, 2004. 50. K. Goda, M. S. Sreekala, A. Gomes, T. Kaji and J. Ohgi, Improvement of plant based natural fibers for toughening green composites—Effect of load application during mercerization of ramie fibers, Composites: Part A 37, 2213–2220, 2006. 51. Y. Nishiyama and T. Okano, Morphological changes of ramie fiber during mercerization, J. Wood Sci. 44, 310–313, 1998. 52. K. Takabe. Highly ordered structure of cellulose. in: Dictionary of cellulose, The Cellulose Society of Japan (Ed.), pp. 102–111, Asakura-shoten, 2000. 53. N. Suizu, T. Uno, K. Goda and J. Ohgi, Tensile and impact properties of fully green composites reinforced with mercerized ramie fibers, J. Mat. Sci. 44, 2477–2482, 2009. 54. M. A. Kabir, M. M. Huque, M. R. Islam and A. K. Bledzki, Mechanical properties of jute fiber reinforced polypropylene composite: Effect of chemical treatment by benzenediazonium salt in alkaline medium, BioResources 5, 1618–1625, 2010. 55. X. Xia, W. Liu, L. Zhou, H. Liu, S. He and C. Zhu, Study on flax fiber toughened poly (lactic acid) composites, J. Appl. Polym. Sci. 132, APP 42573, 2015. 56. A. J. Kinloch, A. C. Taylor, M. Techapaitoon, W. S. Teo and S. Sprenger, Tough, natural-fibre composites based upon epoxy matrices, J. Mater. Sci. 50, 6947– 6960, 2015. 57. L. Yan and N. Chouw, Crashworthiness characteristics of flax fibre reinforced epoxy tubes for energy absorption application, Mater. Design 51, 629–640, 2013.

10 Cellulose Reinforced Green Foams Jasmina Obradovic, Carl Lange, Jan Gustafsson and Pedro Fardim* Department of Chemical Engineering, University of Leuven, Leuven, Belgium Laboratory of Fibre and Cellulose Technology, Åbo Akademi University, Finland Department of Chemical Engineering, University of Leuven, Leuven (Heverlee), Belgium

Abstract This chapter introduces the science, various technologies used, and applications related to the use of biopolymers and biomaterials in the development of porous structures. The main focus is placed on the bio-based foams incorporated with cellulose fibres. The chapter reviews the composition of bio-based foams, processing methods, properties of these porous materials as well as performance and applications of the resulting foams. One section is dedicated to  interesting platform for functionalization of cellulose fibres with layered double hydroxides (LDH). An engineering process for the in situ synthesis of Mg-Al LDH s with pulp fibres is presented as well. LDHs have a particularly interesting property that neither of the constituent precursors have in themselves. LDHs carry a net positive charge due to trivalent aluminium. To counter that charge build up there are intercalated anionic components in between its lamella. Cellulose fibres are naturally acidic facilitating electrostatic interaction with layer double hydroxide in neutral or alkaline medium. Therefore, the synthesis of LDHs, in situ, with cellulose fibres may bring additional benefit to the foam process in which the addition of additives such as flame retardants after the foam has been formed is difficult unless it is inherently part of the matrix. Keywords: Adsorption, wood pulp, cellulose, composite, fibre foam, flame retardant, hydrotalcite, layered double hydroxide, mineralisation, surface engineering, bio-based, starch PLA, vegetable oil

*Corresponding author: [email protected] Anil Netravali (ed.) Advanced Green Composites, (247–274) © 2018 Scrivener Publishing LLC

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248 Advanced Green Composites

10.1 Introduction Polymer foams were first made in the 1930’s and 40’s, with foamed polystyrene being the first polymer foam in 1931. Polyurethane was first invented by Otto Bayer [1] at the beginning of the second world war. It was first used as a replacement for rubber and also as a coating to protect other common materials such as wood and metals. Mass production of polyurethane foams started in the late 1950s. However, long before that, in prehistoric ages, man had mastered the attractive attributes of natural porous materials, like wood and bone, not only for hunting and building, but also for manufacturing various tools, exploiting their advantageous such as specific strength or their thermal insulation properties [2]. Foam is defined as a dispersion of a gas in a liquid or in a solid state. Technical polymer foams are often made by mixing together the solid and the gas phases. The resulting foam has either fully trapped individual air bubbles in it, known as a closed-cell structure or the material may incorporate air tunnels connected to each other which are known as open-cell structures. Foams that have an open-cell structure, like many polyurethane formulations in mattresses for example, are generally more flexible, while closed-cell foams such as styrofoams, are typically rigid. An example of polyurethane foam is shown in Figure 10.1. Foams are never thermodynamically stable. They can only be kinetically trapped within the structure [3]. The gas that is used to create the bubbles in the foam is called a blowing agent. The blowing agent can be either chemical or physical. Chemical blowing agents, when heated to form foams, take part in a reaction or decompose and in the process give off gases which get trapped in the polymeric matrix. Different chemical blowing agents have different

Figure 10.1 Image of a polyurethane foam [4].

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decomposition temperatures depending on their chemical structure, and specific chemical blowing agents can be selected for a particular application. Physical blowing agents are gases that do not react chemically in the foaming process and are, therefore, inert to the polymer forming the matrix. Polymer foams, just like the composite matrices can be divided into either thermosets or thermoplastics, which are further divided into flexible or rigid porous materials. The thermoplastics can usually be broken down or melted and recycled, while thermosets are almost impossible to recycle because the matrix polymers are cross-linked and the structure breaks down irreversibly when heated [5]. Porous polymeric materials play an important role in a vast number of applications. Composed of gas pores integrated into a continuous polymer phase, the form materials whose properties are much more than the simple sum of their parts. Their properties are governed by the structural parameters of the material, such as the average pore size, the pore size distribution, the overall gas-to-volume ratio, and whether the pores are open or closed. The general relationship between these parameters and the overall material properties are quite well understood by now. Solid foams are mostly used as insulators, since trapped air bubbles have low thermal conductivity. The characteristic light weight of foams makes them attractive as filling material in stabilizing hollow structures. Ceramic foams are used in electronic applications, while metallic ones are suitable for aerospace and automotive industry where reducing the weight of every part without reducing the functionality is critical.

10.2 Bio-Based Foams Currently, the polymer foam industry is facing a multitude of issues regarding waste disposal, recyclability, flammability and the effects of blowing agents on the environment. The restriction on the use of chlorofluorocarbons has become very important driving force in making polymer foams more environmentally friendly. At the same time advances in biodegradable foam materials are helping to improve the recyclability and solve issues related to waste disposal. Biodegradable foams have also been developed. While they are water soluble and sensitive to humidity, they are important in reducing the amount of carbon dioxide emission. Bio-based polymers have attracted considerable interest from both academic and industrial researchers as a way to tackle sustainability and environmental concerns. In reality, the search for sustainable development based on renewable bio-based feedstocks has become a major research

250 Advanced Green Composites area. Biomass from plant derived resources is a renewable raw material and with further research, should be able to provide a broad variety of starting materials of monomers and polymers in the future that can be used in foam manufacturing. Bio-based foams are fast growing industry, with large contribution from plant-oil derived foams in construction sector, and from starch foams in packaging area. While the conventional foam industry based on petroleum derived polymers is facing a few issues, the bio-based foams have exhibited similar properties to some of those conventional foams making them fit to go on the market [6].

10.2.1 Starch-Based Foams The starch-based foams have emerged as polystyrene replacements and the development in this field is actually aimed to improve the drawbacks of the starch itself. Starch is universally known as a brittle and water sensitive material. Several approaches have been investigated aiming to overcome these problems. Plasticizer is commonly mixed with starch to improve the foam flexibility. Addition of plasticizer also enables melting of starch below its decomposition temperature, leading to a noticeable improvement in starch processability [7]. Furthermore, in order to enhance water resistance and strength of the foams, chemically modified starches as well as various biodegradable polymers have also been studied. For example, poly(lactic acid), poly(ε-caprolactone), poly(vinyl alcohol), chitosan, poly(ester amide) have been blended with various starches [8, 11]. Another approach is to add reinforcing agents into the starch matrix. This line of action has proved to be an effective method to obtain high-performance starch-based materials [12, 15]. A number of cellulosic fibres such as jute, flax, aspen, corn, soft wood, eucalyptus cellulose, kraft pulp, sugarcane bagasse, cotton, and wheat bran have been used to reinforce starch foams, all being sustainable and fully biodegradable sources. These are presented in Table 10.1. Starch-based foams have been produced by many techniques as shown in Figure 10.2 [10, 26–28]. Each technique has its own processing parameters which affect the product properties including shapes, pore structures and size distribution, density, and mechanical properties. Therefore, the starch-based foam products can be designed to suit a particular application via appropriate processing techniques and adjusting the process parameters. Starch can be melt-processed by adding water or other hydrophilic plasticizers in extruders similar to the conventional polymers. Extrusion has been used in the last few decades for manufacturing low-density foam

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Table 10.1 Starch-based foams reinforced with various fibres. Fibre

Starch

Process

Reference

Aspen fibres

Native corn starch Baking process

Lawton, et al. [6]

Kraft fibres

Cassava starch

Baking process

Kaisangsri, et al. [7]

Aspen fibres

Potato starch

Baking process

Shogren, et al. [8]

Cotton linter Potato starch fibres, Hemp fibres, Cellulose fibres, Sugarcane fibres, Coconut fibres

Twin screw extrusion

Bergeret and Benezet [9]

Wheat straw fibres, Hemp fibres, Cotton linter fibres, Cellulose fibres

Potato starch

Twin screw extrusion

Benezet, et al. [10]

Sugarcane fibres

Cassava starch

Single screw extrusion

Debiagi, et al. [11]

Bleached softwood fibres

Wheat starch

Baking process

Glenn, et al. [12]

Wheat bran

Starch

Twin screw extruder

Robin, et al. [13]

Eucalypt fibres

Cassava starch

Baking process

Salgado, et al. [14]

Eucalypt fibres

Cassava starch

Thermopressing Schmidt and process Laurindo [15]

Jute fibres, Flex fibres

Tapioca starch

Baking process

Malt bagasse

Cassava starch

Thermopressing Mello and Mali process

Kraft fibres

Cassava starch

Moulding

Cassava fibres

Cassava starch

Thermopressing Carr, et al. [24] process

Soykeabkaew, et al. [16]

Kaisangsri, et al. [23]

252 Advanced Green Composites

1 mm (a)

1 mm (b)

30 mm 1 mm (d)

(c)

500 um

500 m (e)

(f)

Figure 10.2 Starch-based foams prepared by (a) extrusion, (b) hot mould baking, (c and d) microwave heating, (e) freeze-drying, and (f) solvent exchange. [28]

materials with physical blowing agents, by using a variety of polymers including polystyrene, polyethylene, polypropylene and poly(lactic acid). Generally, extrusion processing system produces the foam with large pore size and that leads to the production of low density products. This can be useful for the application where reducing the weight is a high priority. In applications that require the well moulded shape such as disposal containers, baking compression process may be well suited as the foam is expanded within a closed mould. Starch-based foams with various shapes such as cups and trays can be made by baking a mixture of starch and water in a closed and heated mould for several minutes. During baking, the starch granules gelatinize into a thick paste while the water rapidly boils, causing the paste to expand. When the starch paste fills up the mould, residual water further evaporates. The starch foam is then dried and takes up the

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shape of the desired mould. Microwaved foams allow thick cell walls which can be suitable for the application where strength and stiffness are essential. On the other hand, applications which require micro-scale and uniform pores with interconnectivity such as scaffold or drug delivery system, the freeze drying and solvent exchange processes might be more suitable. A recent supercritical fluid extrusion technology seems to be more versatile and controllable, thus, a wide range of foam structures with desired properties can be designed [28]. Numerous cellulose fibres have been incorporated in starch-based systems forming composite foams, which have generally provided a significant improvement in mechanical properties of the materials at appropriate fibre content. These composite foams have improved mechanical properties since both starch and cellulose fibres are similar in their chemical composition. Therefore, good compatibility and strong interaction between them are established, leading to satisfactory stress transfer in the composite system. Generally, it has been observed that longer the fibres added to the system, the stronger and stiffer the foams are. Moreover, as expected, with increasing fibre contents, the strain or flexibility of the composite foams decreases and the foam density increases [28].

10.2.2 Foams Based on Vegetable Oils Vegetable oils and fats are triglycerides, i.e., esters of glycerine and different fatty acids. Triglycerides are excellent renewable raw materials for a range of products including polymers. Vegetable oil-based polymers are gaining popularity due to some attractive properties related to the specific structure of the oils. They are also inexpensive and available worldwide. Historically, vegetable oils have been used for air drying coatings. Reactive sites in all oils are ester bonds and double bonds, and in a few oils other groups such as hydroxyl or epoxy may also be available. Polyurethanes are prepared by reacting polyol components with isocyanates. At this time, it is realistic to make only the polyol component from vegetable oils for foam applications [29]. Vegetable oil based polyols for rigid and flexible foams can be prepared through direct oxidation of oils, epoxidation followed by ring opening, hydroformylation, ozonolysis, and transesterification. Commercial and academic research is most focused on polyols derived from soybean oil [30–32] or soybean oil derivatives [33]. Nevertheless, there are reports of successful bio-polyol synthesis from corn oil [34], castor oil [35], sunflower oil [36], but rarely from rapeseed oil [37], linseed oil [38], and tung oil [39]. Vast combinations of fibres and additives have been

254 Advanced Green Composites explored as reinforcement for vegetable oil-derived foams as presented in Table 10.2. Long fibre reinforcement is aimed at altering or improving the mechanical performance of a system, whereas short fibre reinforcement holds advantages in both thermal and mechanical properties. Cellulose fibres have many benefits when used in foams derived from vegetable oil, beyond their basic advantages of renewability and potential biodegradability. Moisture content of cellulose fibres allows improved foaming kinetics and enhances interfacial interactions between fibres and foam matrix [50]. Moisture in cellulose fibres assist in the blowing process, where water is necessary, resulting in foams with slightly higher cell densities. For instance, flax and hemp fibres cause noticeable improvement in mechanical properties compared to vegetable-oil-based foam without any reinforcement [51, 52]. Therefore, the incorporation of cellulose fibres holds potential towards improvements in a different set of foam properties.

Table 10.2 Vegetable oil-derived foams reinforced with various fibres. Fibre

Polyol

Process

Reference

Pine fibres

Castor oil

Moulding process

Aranguren, et al. [40]

Bleached Kraft fibres

Soy oil

Moulding process

Banik and Sain [41]

Flax fibres

Rapeseed oil

Hemp fibres

Petrochemical compound

Free-rise method

Kuranska and Prociak [42]

Pulp fibres

Lignin

Free-rise method

Xue, et al. [43]

Aspen fibres

Soy oil

Free-rise method

Khazabi, et al. [44]

Wood flour

Castor oil

Free-rising in a mould

Mosiewicki, et al. [45]

Wood fibres

Tall oil

Free-rise method

Cabulis, et al. [46]

Refined wood fibres

Soy oil

Free-rise method

Zhu, et al. [47]

Wood fibre

Soy oil

Free-rise method

Gu and Sain [48]

Wood fibres

Soy oil

Free-rise method

Chang, et al. [49]

Hemp fibres

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10.2.3 Foams Based on Poly(Lactic Acid) Poly(lactic acid), (PLA), is a biodegradable and compostable polyester belonging to the aliphatic polyester family derived from α-hydroxy acids. It can be produced through ring-opening polymerization of lactic acid [53]. PLA is a rather brittle thermoplastic with high strength and modulus. It is derived from renewable resources such as corn, sugar beets and sugarcanes, or even wheat and other starch-rich products [54]. Considering PLA’s processing costs and mechanical properties, this environmentally friendly biopolymer is considered as a promising replacement for polystyrene foam products in such commodity applications as packaging, cushioning, construction, thermal and sound insulation, and plastic utensils [55, 56]. PLA foaming has mostly been conducted by dissolving a physical blowing agent in the PLA matrix. Cell nucleation and growth occurs through the thermodynamic instability generated from the super-saturation of the blowing agent. The foam structure is then produced by expelling the dissolved gas from the PLA/gas mixture. Foam products are made with cell stabilization as the temperature reaches below the PLA’s glass transition temperature, which is around 60 °C [57]. In order to manufacture fully degradable ‘green’ foams with improved mechanical properties, a couple of researchers have investigated PLA foaming behavior using a few cellulose fibres and forming fibre-reinforced composite foams. These results are presented in Table  10.3. Silane was occasionally used as a coupling agent to create a strong interface bonding between the PLA and the fibres. The foaming results showed that having fibres improved both the cell density and the foam’s crystallinity. Cellulose fibre reinforced PLA foams can be manufactured by supercritical carbon dioxide assisted foam extrusion process, the same method that could be easily scaled-up towards real industrial applications. Literature research shows that the addition of cellulose fibres provides a wider processing window for PLA foam production. Based on the lower applicable temperature profile, a greater degree of crystallinity and improved melt strength could be achieved. It was demonstrated that cellulose fibres increase the melt viscosity and promote the heterogeneous cell nucleation, obtaining high porosity foam structures. However, due to the fibre distribution and the weak fibre-PLA adhesion, the fibre reinforced composite foams have less uniform cell structure and increased open cell ratio compared to the non-reinforced foam. It is believed that further chemical or physical modifications, such as reinforcement and flame retardancy, could promote the market penetration of PLA foams in technical application fields as well [58].

256 Advanced Green Composites Table 10.3 PLA-derived foams reinforced with various fibres. Fibre

Blowing agent

Process

Reference

Softwood kraft fibres Black spruce fibres

N2

Injection Moulding process

Ding, et al. [59]

Cellulose fibres Basalt fibres

Supercritical CO2

Twin screw extruder

Bocz, et al. [58]

Wood flour

CO2

Batch process

Matuana and Faruk [60]

Flax fibres

N2

Injection Moulding process

Pilla, et al. [61]

Recycled paper fibres

N2

Injection Moulding process

Kramschuster, et al. [62]

Microcrystalline cellulose

CO2

Batch process

Boissarda, et al. [63]

10.3

Surface Engineering of Cellulose Fibres Used in Foams

Wood is a natural composite. The main characteristics common to all wood fibres are the intercellular middle lamella that glues the fibres together, the mesh or cage like primary layer and two or three structurally important secondary layers. The dimensions of the fibres, the thickness and structure of the cell walls and their chemical composition differ from species to species and vary within the wood according to the location in the stem, growth seasons, environment, geographical location and the possible interruptions in the normal growth [64]. The three major components in wood fibres, namely, lignin, cellulose and hemicelluloses are distributed throughout the cell wall structure in different proportions. Cellulose serves the wood by being responsible for the microstructural property of different cell wall layers and it is the major contributor to the overall physical features of the fibres. Cellulose molecules arrange regularly, gather into bundles, and determine the framework of the cell wall. Cellulose is a natural high molecular weight linear polymer composed of glucose monomers,

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with cellobiose as the basic repeat unit. Cellobiose consists of two glucose units combined together and has eight free  alcohol  (hydroxyl) groups (OH), one acetal linkage and one hemiacetal linkage, which give rise to strong inter- and intra-molecular hydrogen bonds. The poor mechanical properties of bio-based matrix materials in foam manufacturing compared to conventional petroleum-based polymers can be improved by adding cellulose fibres. For reinforcement of bio-based matrix with cellulose fibres, several problems occur along the interface due to the presence of hydrophilic hydroxyl groups on the fibres. This hydrophilic nature hinders the effective reaction with the bio-based matrix. In addition to this, pectin and waxy substances cover the reactive functional groups of the fibre and act as a barrier to interlock or chemically bond with the matrix. To enhance the effectiveness of interfacial bonding, fibre surface needs to be modified. Functionalization in chemistry refers to a molecular level synthesis, obtaining a specific chemical group or a material characteristic and to the practicality of the material that is designed for a certain application. As we are dealing with cellulose fibres in this chapter, the reference is often made to a molecular level surface modification. Common examples include grafting of co-polymers [65], polyelectrolyte adsorption [66, 67], plasma gas treatment [68, 69] and enzymatic treatments [70, 71]. Apart from molecular level functionalization, inorganic particles in different size ranges can be used for surface modification as well [72–74].

10.3.1 Chemical Modifications of Cellulose Fibres As already mentioned, the main obstacles in the use of cellulose fibres in foam production have been the insufficient compatibility between the fibres and the foam matrix, and the inherent high moisture absorption by fibres, which could result in dimensional changes of the fibres that may lead to degradation of mechanical properties of the foam. Cellulose fibres are amenable to chemical modification due to the presence of large number of hydroxyl groups on the surface of the polymer. These hydroxyl groups may be activated or new moieties can be introduced with proper chemical treatment of the fibre. Chemical modifications employ chemical agents to modify the surface of the fibres or the whole fibre throughout. The chemical treatments of cellulose fibres are intended to: Improve the adhesion between the fibre surface and the polymer matrix Increase the fibre strength

258 Advanced Green Composites Increase moisture resistance of cellulose fibre (making it hydrophobic), and Improve mechanical properties of the material Chemical treatments provide important and effective means to remove non-cellulosic components in cellulose fibres and add functional groups to enable better bonding to the polymer. Many chemicals have been screened in laboratory experiments for potential to enhance fibre/matrix interface, such as sodium hydroxide [75], peroxide [76], organic and inorganic acids [77], silane [78], anhydrides [79] and acrylic monomers [80]. Synthetic layered double hydroxide (LDH) appears as a promising platform to modify cellulose fibre surfaces. It provides cellulose fibres with ampholytic character and reduces liberated heat under forced burning. LDH is also able to mineralize wood fibres if the synthesis route is correctly chosen [81]. These characteristics can be further exploited in diverse applications where fibres require ability to absorb both cationic and anionic substances. On the other hand, flammability issues are important whenever the fibres are used in indoor applications. Pulp fibre industry can readily apply LDH synthesis into the existing processes and take advantage of the proposed hybridization. LDH may also be useful in more general terms in fibre technology, since LDH formation does not require complex chemistry.

10.3.2

In Situ Synthesis of Hybrid Fibres

The process that involves stabilization of crystal formation by a metal cation is well known and widely exploited method in LDH preparation. It has been proven that during the co-precipitation step not only inorganic anions but also graphene [82], pre-polymers [83], and biologically active molecules [84] may be incorporated in LDH structure. The precipitation strategy depends on the preferred hybrid structure. The precursors should be chosen in order to promote affinity towards the intercalated anion. Usually nitrates are applied in this respect. Also, the intercalated anion is often dissolved in excess from the calculated fraction of trivalent cations in the crystalline lattice. Widely, the co-precipitation is achieved either at high super saturated (hss) or at low super saturated (lss) conditions [85]. The hss synthesis route may be chosen if the material is allowed to age hydrothermally. The problem in hss is in the continuous pH change during the particle nucleation. Buffers, however, can be applied. The lss system is preferred if charge density ratio needs careful control. Neither one of these synthesis routes is able to produce particles with narrow size range

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distribution, wherefore, a rapid nucleation procedure has been developed [85]. Subsequent treatment after the LDH formation, especially if the material is to be used as a catalyst or if the intercalated anion exchange is required, involves calcification process. Catalysts may be used in any temperatures above approximately 350 °C while in the case of ion exchange the temperature should be kept below approximately 525 °C. The limit is set in the latter case by the LDH’s rehydration ability, i.e. the so called memory effect first explained by Miyata [86]. The crystallographic data revealed that, once calcified, the mineral acquired its original structure upon hydration in aqueous environment. Applying the memory effect procedure allows the LDH to incorporate many different anions without using high ionic concentrations or elevated pH. Application of inorganic particles as a platform for pulp fibre functionalization has received little attention. Considering LDH as a tool for pulp fibre functionalization, only a handful of scientific articles and patents can be found. For example, precipitated LDHs have been utilized as fluoride scavenger in aqueous media [87], as a flame retardant in paper [88], in paper filler and fluorescent whitening agent formulation [89], catalyst in lignin depolymerisation [90]. Effective LDH particle nucleation on fibre surface can be successfully conducted in aqueous medium of lss, hss and slow urea hydrolysis (Uhyd) synthesis methods. In these synthesis procedures the applied pulp fibres were acting as templates for nanometer (70 nm), sub-micrometer (200 nm) and micrometer (2 μm) size LDH particles, respectively. In situ nucleation of LDH from hss system with fully bleached Kraft pulp fibres in an autoclave does not induce particle synthesis or migration into the fibre wall, even if the fibres are fully saturated and swollen in alkaline medium (pH = 10) and elevated temperature [91]. The results of synthesis via lss route indicate slight particle migration into the fibre wall. Considerable particle nucleation inside the fibre wall was discovered with transmission electron microscopy after hydrothermally induced Uhyd. LDHs that were bound to fibre’s external surface had a marginal effect on the fibre flexibility. Synthesis conditions were found to induce cellulose depolymerisation reactions. LDH acts as a catalyst. The cellulose degree of polymerisation (DP) followed each synthesis in the order: hss > lss > Uhyd, suggesting that Uhyd had the greatest effect. Utilizing sequenced in situ particle synthesis via Uhyd and homogeneous co-precipitation results high mineral content on the fibre surface. Fine structures are prone to act as LDH nucleation substrate and favored over fibres. LDH provide a relatively dense structure enveloping of the fibre outer cell wall. Synthesis of LDH at alkaline

260 Advanced Green Composites conditions liberates organic substances from the thermo-mechanical pulp fibres. Some of the compounds are re-deposited on the LDH surface. Significant reduction in specific exothermic heat was observed after each synthesis route. Reduction in cellulose molecular length or DP influenced the heat of combustion during the first exothermic reaction at around 320 °C. Influence of LDH in enthalpy reduction was weight related. Carboxylic acids are created in higher relative amounts during the slow combustion in comparison to rapid temperature increase that produces diones, most likely due to simultaneous removal of water and hydroxyls from the LDH. The effect is amplified by the particles present in the fibre wall. The LDH particles that were synthesized via lss method had the highest capacity for sulphate containing probe molecules while the better crystallized particles from hss and Uhyd had the highest affinity. Most of the acidic groups in cellulose pulp fibres were free after LDH synthesis [91].

10.3.2.1 Topology and Particle Content on Hybrid Fibres The nano-sized LDH particles were densely packed on the fibre’s surface and the topology of the functionalized fibres differed as shown in Figure  10.3. However, clear distinction in between individual LDH particles was difficult to make. In the freeze dried pulp fibres the particles appear to be coarse, growing as if facing the brucite layers in right angles with respect to the fibre’s surface. It has been proposed that the orientation of LDH particles may be influenced by the substrate and especially by the hydroxyl groups that it contains [92]. The surfactant itself did not induce any additional topological changes in the applied range (6% w/w).

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Figure 10.3 Fibres morphological changes upon modification via lss route using chloride salts. The reference fibres (a) show clearly visible S1 layer. Spruce fibre’s surface topology changes greatly at submicron level after modification with high number of LDH (100 mM) (b). Arrows are pointing towards the fibre axis. [93]

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Fibres were characterized by TEM instrumentation that revealed LDH particles reside mainly on the fibre’s surface and lumen as well as on fine structures as can be seen in Figure 10.4. Interestingly, the fibre cell wall was void of LDH particles. It appeared that the sequential synthesis of LDH particles on bleached thermo-mechanical pulp fibres effectively fibrillated the fibres by forming thread like structures. Energy dispersive spectra (EDX) showed small amount of sulphur on the surface of the foamed reference panel but not in

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Figure 10.4 TEM images of reference bleached thermomechanical pulp fibres (a) and hybridized fibres (b, c). LDH particles appear to have nucleated mainly on fibre surface and lumen as well as to fine structures. [93]

262 Advanced Green Composites the middle layer of the panel. Surfactant that was used as a foaming agent must have been adsorbed onto the fibres during the dewatering of foamlaid fibre web. These webs, when prepared from the hybridized fibres, contained sulphur throughout the cross-section. Estimation for the effective absorption of sodium dodecyl sulfate (SDS) surfactant on LDH containing fibres can be made via aluminium to sulphur ratio. It was noted that the Mg2+ to Al3+ ratio in panel prepared from hybridized fibres did not change markedly while Al3+ to S6+ was slightly higher in the middle layer. The surfactant was prone to adsorb onto the panel surface similar to reference foamed structure.

10.3.2.2 Foam Formation A sheet former can be used to prepare foamed structures with average grammage of 100 g·m 2 [43]. SDS was used as a foaming agent in this case. The aqueous fibre sludge was vigorously mixed with prefabricated foam. The amount of SDS was optimized to 0.15–0.2 g·L 1 leading to 60–70% air content in the foam. Stabilized foam was decanted into the sheet mould. The system is designed to allow fibres to be oriented with the suspension flow. Settling of foam is followed by vacuum suction. Detached sheet is dried on a separate suction table that is equipped with a 5 mm wide slit enabling for air to pass through the foam-laid sheet [81].

10.3.2.3

Combustion Behavior of Foams

The outcome of a combustion test in a cone calorimeter is influenced, for example, by the vertical distance of the sample from the external heat source, applied heat flux, sample holder system and the physical dimensions of the specimen. Applicability of cone calorimeter in prediction of hazardous fires has been discussed by others [94]. In combustion experiment the distance of the sample surface from the external heat source was 60  mm. This distance is relatively long and may cause the centre of the sample to absorb higher relative irradiance per unit time than the sample edges [95]. However, the tested low weight fibre foams were dimensionally unstable under forced burning. The relatively low irradiance (25 kW∙m 2) was chosen based on that characteristic and was targeted to address the ignition phenomenon. Sample thickness is also an important factor in cone calorimeter experiments. Foam-laid fibre panels are meant to be applied as a bulk material, wherefore, the chosen 7.5 mm thickness is within a practical limit. Typically, several such layers are combined to a lamellar structure that is eventually compressed to appropriate density in a suitable mould.

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Thinner samples tend to show higher peak heat release rate values while thicker samples burn longer allowing better estimation of flame propagation within the structure. Also, the cut samples were placed on a 20 μm thick aluminium tray. Aluminium is a heat conductor and transports irradiated heat away from the sample. The effect of heat transport is assumed to be small, however, because the wood fibres are good heat insulators, and, because the contact interface area in between the sample and aluminium tray was considerably smaller than the geometrically measured flat surface area of the specimen. Although cone calorimeter cannot mimic real fire due to measurement restrictions, its importance is in data related to the material behavior after heat radiation impinges the sample surface. In cone calorimeter the heat release rate is the most important factor and is calculated from the flue gas oxygen consumption. Anything that limits oxygen usage upon flaming, such as generation of soot, decarbonation, temperature drop within the material, water evaporation, etc., will affect the observed heat release rate [96]. The in situ synthesized LDH particles on fibres provided reduction in peak heat release rate, CO2 production rate and smoke yield as shown in Figure 10.5. Tailing in heat release rate signal predicts that both systems are charring to some extent. The mass loss rate is reduced in LDH-fibre foam sample indicating slower material volatilization. Narrow heat release rate profile and rapid combustion was expected due to the sample thickness. Carbon monoxide (CO) yield and production rate increased relative to the fibre foam reference. Partial oxidation that appeared before the advancing heat release in both samples, also related to CO production and material pyrolysis at the reaction zone front, appeared similar in rate and yield. The production rate of CO in material glowing phase, i.e. after the combustible volatiles have been released but the temperature remains sufficiently high to maintain the material pyrolysis, however, was approximately three times faster in case of LDH-fibre foam. Combustion efficiency, yet another important characteristic in cone calorimeter experimentation, defined as a ratio of total heat evolved to mass loss change until flame out, was reduced by LDH. Some researchers have suggested that materials safety in case of fire can be evaluated and characterized indicatively from cone calorimeter results by simple calculations based on peak heat release rate, time to ignition and total energy release [97]. For example, the so called flash over propensity, defined as the ratio of peak heat release rate to time to ignition drops from initial 22 (high risk) of fibre foam to 6 (moderate risk) in LDH-fibre foam. Total energy release provided a low risk assessment for those samples. Also,

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Figure 10.5 Mass release rate of CO2 (a, b) and CO (c, d) (black lines) during the oxidative pyrolysis compared against heat release rate (blue line) in fibre foam reference (a, c) and LDH-fibre foam(B,D), and, data from mass fraction extent (e) and combustion efficiency (f) are presented. Time to ignition (tIG) and time to flame out (tFO) are marked with red vertical lines. [93]

the smoke production subsided within 15 s after ignition in case of LDHfibre foam while, for the reference foam, it took nearly 22 s. Amount of soot was about four times greater in reference fibre foam sample. The mineral containing foam will therefore produce less fine particulates upon fire. In comparison, for example, the ammonium polyphosphate treated plywood released significantly higher amounts of smoke than the reference sample while the peak heat release rate values and the relative amounts of CO formation were similar to reported results [98]. It was interesting to note that the flame subsided slower in LDH-fibre foam; albeit it did not develop as it did for the reference fibre foam. The LDH particles seem to act as

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heat barrier and reduce the liberation of burning volatiles. An example of material morphology after combustion is given in Figure 10.6. Carbonized organic substances are seen on reference fibre foam while particles that resemble features of fibres are seen in LDH-fibre foam.

10.4 Prospects At the beginning of the 21st Century and unlike any other class of materials, bio-based porous materials with an attractive balance of low cost and high performance are essential for meeting the demands of the growing global population with respect to health, buildings, insulators, food and energy. Renewable materials play a key role in sustainable development. Development of new commercial products or new technologies is the objective of continued research in both academic and industrial fields. The production of renewable and green bio-based foams has been demonstrated to be technologically feasible and capable of substituting fossil feedstocks. Renewable vegetable oils and bio-based polymers have been proposed as composite matrices/resins reinforced with cellulose fibres in foam production, and this trend is continuing to grow with promising results. Some of the vegetable oil based polymers and precursors are currently commercially viable and there are companies that offer derived biopolymers for certain applications. Bio-based materials continue to provide new opportunities in some particular realms, including biomedical applications, in the near future. It is inevitable that in the fast approaching future, renewable feedstocks will be the only alternative to finite petroleum feedstock reserves. The initial use of biopolymers and biomaterials was presented as a partial replacement of synthetic polymers, and little more than that. However, the current

266 Advanced Green Composites trend is to increase the percentage of bio-based materials, maintaining good overall performance, and/or developing specific properties.

10.5 Summary Polymers are among the distinctive features of 20th century science, and the intense outgrowth and tremendous importance of polymeric materials is a testament to their amazing versatility and unique properties. Foamed polymer materials represent an important extension of polymer spectrum with enhanced properties with significant weight (density) reduction. With their uses ranging from insulators to automotive applications and from mattresses to medical purposes, polymeric foams pervade all areas of our everyday lives. If this growth is to continue, especially taking into account strict environmental regulations, a strong understanding of the basic science and the major technologies behind foam production is critical. The development of different technologies for polymer foam manufacturing and newly designed processing equipment have been the key factors that promoted the development of polymer foams. The productions of both thermosetting and thermoplastic polymer materials have contributed to the development of foam industry in the last few decades. The foaming methodology consists of introduction of gaseous phase into a liquid and subsequently solidifying the liquid before gaseous bubbles collapse. Thermoplastic and thermoset foams are produced using very different processes. Thermoplastic foams have unique thermal reversible character and are strongly dependent on the kind of equipment used in manufacturing process. On the other hand, the production of thermoset foam seems to depend entirely on the reaction kinetics and material formulation. The widely used renewable materials are derived from wood and based on cellulose. However, major efforts are made to identify alternative uses of wood and agricultural resources. A wide variety of monomers or building blocks can be obtained from biomass. These can be extracted using chemical or biotechnological routes. Materials derived from biomass can be used in production of different type of foams using existing equipment and technology used for conventional materials. On the other hand, these foams have to be well performing in order to compete with highly developed currently used materials. In general, the more diverse functional groups of bio-based material give the resin unique possibilities to tailor the properties of final product. It has been observed that lignocellulosic fibres are able to provide various beneficial characteristics to the polymer foam systems. As a result, these cellulose fibres are interesting candidates due

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to the reinforcing effect they can bring to foams. It is believed that further chemical or physical modifications of cellulose fibres, such as reinforcement and flame retardancy could promote the market penetration of biobased foams in technical application.

Acknowledgements The author would like to express his gratitude to the many people who saw him through this book chapter; to all those who provided support, talked things over, read, wrote, offered comments, and assisted in the editing, proofreading and design.

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11 Fire Retardants from Renewable Resources Zhiyu Xia1, Weeradech Kiratitanavit1, Shiran Yu1, Jayant Kumar2, Ravi Mosurkal3 and Ramaswamy Nagarajan1* 1

Department of Plastics Engineering and Center for Advanced Materials, University of Massachusetts Lowell 2 Department of Physics and Applied Physics and Center for Advanced Materials, University of Massachusetts Lowell 3 Bio-Science & Technology Team, Materials Science and Engineering Branch, US Army Natick Soldier RDEC, Natick, MA, USA

Abstract Rapid expansion in utilization of polymers for construction, transportation, electronics and protective textile applications has necessitated the development of new flame retardants (FR) to ensure fire safety of these products. The toxicity and environmental persistence of certain classes of commercially available halogenated FR has led to the quest for safer and ‘greener’ alternative FR preferably obtained from renewable feedstock. This chapter attempts to provide a brief overview of the most promising bio-based FR alternatives and their mechanism of action. Biobased alternatives with the potential for imparting FR characteristics are classified and described based on their elemental composition (nitrogen and phosphorus containing compounds), chemical class (carbohydrates, nucleic acids, proteins, polyphenols) and synergistic combination of one or more of these with inorganic materials such as nanoclay. The bio-based additives typically facilitate in lowering the rate of flame propagation, increasing char formation and reducing the total heat released upon burning. The mechanisms of action of these additives appear to be predominantly through formation of char that can provide both thermal and gas barrier, insulating and protecting the polymer substrate and reducing the flammability. Based on recent and unexpected results from blends and coatings of polyphenols with polyamides and polyurethanes, there is increasing speculation for gas phase radical scavenging action from polyphenols acting in synergy with char formation in these systems, providing self-extinguishing behavior. Finally, the limitations of natural materials including consistency of feedstock as well as

*Corresponding author: [email protected] Anil Netravali (ed.) Advanced Green Composites, (275–320) © 2018 Scrivener Publishing LLC

275

276 Advanced Green Composites thermo-oxidative stability issues during processing are also discussed. Exploring the use of renewable materials as feedstock for FR alternatives is quite promising. In addition to sustainability aspects, it has unlocked exciting new possibilities in fundamental understanding of fire retardants and their mechanisms of action. Keywords: Fire retardant, renewable resource, sustainability, phosphorus/ nitrogen-containing flame retardant, carbohydrate, intumescent flame retardant, polyphenol, tannin, chemical modification

11.1 Introduction Fire retardants help in reducing fire related injuries, loss of human life and property by lowering the probability of the fire spreading uncontrollably upon accidental exposure of flammable materials to an ignition source. Fire retardants curtail rapid propagation of fire and reduce the ultimate size of the fire while increasing the time available for escape and substantially improve the probability of survival for the occupants. Typically, flame retardant (FR) additives are used to reduce fire propagation in materials that are otherwise highly flammable and meet fire performance criteria (often stipulated by regulations). Over the past decades, fire retardancy of polymers has typically been achieved using synthetic compounds containing halogens, nitrogen, and/or phosphorus. The concerns regarding the continued use of certain types of current fire retardant additives are multifold. Bioaccumulation of certain types of FR due to their diffusion from products containing these additives and their environmental persistence has been well documented [1–5]. On the other hand, the release of toxic combustion products such as corrosive acids, hydrogen cyanide, phosphorus containing compounds, halogenated dibenzodioxins and furans is also of significant concern [6]. The toxicity of combustion products depends on numerous variables including the elemental composition, chemical structure, combustion conditions (fuel/air equivalence ratio) and non-flaming versus flaming behavior [7]. Finally, the availability of feedstock for the production of phosphorus containing FR is an important consideration from a long-term sustainability perspective. However, compromises have been made to solve more pressing problems with materials’ performance and fire safety considered as having higher priority over unpredictable long-term impact on environment and human health. For example, polymeric halogenated FR, such as brominated polystyrene, has been proposed as an alternative to the conventional small-molecule based halogenated FR compounds (hexabromocyclododecane, ‘HBCD’ and decabromodiphenyl ether, ‘DBDPE’) by some of the major halogenated FR manufacturers. However, sufficient

Fire Retardants from Renewable Resources 277 toxicology data is still not available and the impact of these new alternatives on the environment and humans is unknown. With increasing toxicity and environmental concerns, researchers from both the industry and academia have explored the use of naturally occurring feedstock for the synthesis of less/non-toxic FR additives with minimal environmental impact and comparable efficiency. Broadly, the methods for achieving equivalent fire performance involve noncombustible and preferably inert fillers that can be intumescent and release water during endothermic processes, and additives that promote char formation and inhibit combustion of materials or composites. The research efforts are also focused on expanding bio-based options and narrowing the performance gaps between the bio-based alternatives and well established commercially available FR materials that are toxic or environmentally persistent. The control of flammability of polymers using non-toxic and environmentally friendly yet effective additives remains as one of the most important challenges especially in polymers utilized in building/construction, electronics, composites, automobiles and protective fabric applications. With constant expansion of polymers being used in multitude of applications, the quest for safer and effective FR additives that can be derived from renewable resources is becoming increasingly relevant. The established modes of FR action provide general guidelines in terms of the elemental constituents and compositions that are beneficial for the development of safer and eco-friendly FR obtained from renewable feedstock. It is becoming obvious that certain types of natural compounds can be utilized as resources for phosphorus. For example, phytic acid and deoxyribonucleic acid (DNA) contain phosphorus. Casein, hydrophobin and whey protein contain nitrogen and have been shown to be beneficial for FR applications [8–10]. Polysaccharides such as chitosan, alginate and starch are good char forming compounds [11–14]. Coating of natural polyelectrolytes using sequential electrostatic layering techniques (layer-by-layer techniques) on fabric has been reported to provide promising self-extinguishing characteristics [15, 16]. The current understanding of these systems suggests that the bio-based molecules act as intumescent FR by forming char layer that provides a thermal barrier in the condensed phase. Polyphenols represent another class of renewable materials that form carbonaceous char upon burning. They parallel the thermal stability and char forming abilities of phenolic resins. Even though reports on materials with FR characteristics obtained from renewable resources have been prevalent in scientific literature over the past several years, very few have been commercialized. This chapter provides an overview of the current state-of-the-art on fire retardants obtained from renewable resources along with their limitations.

278 Advanced Green Composites Based on the mode of action, FR can be classified into three categories: halogenated hydrocarbon – antimony oxide, phosphorus/nitrogen-based compounds, and metal hydroxides. Synergistic combination of radical quenching effect in the vapor phase provided by halogenated FR along with antimony oxide is well-known to provide universal FR solution for many polymer applications. Metal hydroxides normally act as a heat sink by endothermic degradation into metal oxides accompanied with release of water vapor to dilute gas phase. However, the mechanical properties of the final compounded polymer blends are normally adversely affected by the heavy loading required for the desired FR characteristics. The FR performance at 20–30 wt% loading range is hardly comparable to that achieved by the use of halogenated or phosphorus-based FR. Phosphorus/nitrogen-based FR have also been extensively developed because of the radical quenching mechanism from organophosphorus compounds and intumescent char forming mechanism from polyphosphate/polyol/amine-based combinations. The polyphosphate/polyol/amine-based intumescent FR is normally used as flame resistant coating applied to the surface of materials. The combination of acid source (phosphoric acid), charring agent (polyol) and blowing agent (amine) often renders it both versatile and complex FR system which needs optimized ratios of each component. Phosphorus and nitrogen containing compounds and carbohydrates obtained from renewable resources have been reported to have a role in FR applications. These compounds are categorized based on their source and function in the context of fire retardant behavior and presented in Table 11.1. Comprehensive reviews on the use of biomacromolecules as FR for textiles, covering various phosphorus-rich materials (DNA, phytic acid, casein and hydrophobin), proteins (whey protein and keratin), carbohydrate (chitosan, alginate and starch) and plant extracts (from spinach and banana pseudostem) have been attempted [17, 18]. This chapter summarizes the impact of these bio-based materials on char forming characteristics, heat release characteristics and flame test performance when applied as coatings or melt blended into plastics. A summary of research studies on bio-based FR molecules has been presented in Table 11.2.

11.2 Fire Retardant Additives Based on Phosphorus and Nitrogen from Renewable Resources Phosphorus and nitrogen containing materials have been found to be very effective in FR applications due to multiple mechanisms of action quite often in synergy. Upon combustion, phosphate groups typically degrade

Fire Retardants from Renewable Resources 279 Table 11.1 Bio-derived materials and their roles in imparting FR characteristics. Role in imparting FR characteristics Bio-derived materials

Acid source to catalyze char formation

Carbon source for char formation

Blowing agent

DNA

√

√

√

Phytic Acid

√

√

Casein

√

√

Hydrophobin

√

√

Whey Protein

√

Chitosan

√

Alginate

√

Starch

√

Nanocellulose

√

Tannin

√

Lignin

√

√

√

into phosphoric acid. The released phosphoric acid can catalyze carbonization and form polyphosphoric acid by dehydration. Thermal degradation of polyols is also known to release water and form conjugated structures that cyclize, thus yielding thermally stable char. Nitrogen-containing functional groups/compounds, such as ammonium or melamine, are normally chosen as the cations in the phosphate containing FR molecules. In these cases, ammonia can be released upon combustion followed by formation of polyphosphoric acid along with carbonized char. Ammonia gas also expands the char before it solidifies and fills the char and displaces the air. Bio-macromolecules rich in phosphorus and nitrogen that can provide FR action are described under categories of nucleic acids, proteins and carbohydrates in this section.

11.2.1 Nucleic Acids Deoxyribonucleic acid has recently been identified as a potential green intumescent FR biomacromolecule due to its distinct chemical composition

16% at 700 °C N2 11% at 700 °C Air

15% at 600 °C N2 42% at 700 °C N2 from PCFC

Epoxy/2.5 wt% nanoclay-g-DNA (DNA 0.2 wt%)

Nonwoven PLA/45 wt% Phytic Acid

Cotton/18 wt% 32 bilayers (12 wt% phytic acid)

DNA-nanoclay Phytic Acid

Phytic Acid-Chitosan (LBL) Phytic Acid-Chitosan Polyelectrolyte Complex (PEC)

13% at 850 °C Air

LDPE/40 wt% DNA

259 100 W/g

465 278 W/g

1542 1220 kW/m2

1123 695 W/g

1036 562 kW/m2

PA6/10 wt% DNA

DNA

973 563 kW/m2

2

1300 629 kW/m2

97 57 kW/m2

No

1180 512 kW/m

PP/10 wt% DNA

DNA (surface coating)

13% from cone calorimetry

125 kW/m2 ignition

PET/10 wt% DNA

Cotton/14 wt% 20 bilayers

DNA – Chitosan

35% at 600 °C N2 19% at 600 °C Air

Char residuea

Peak heat release rate (pHRR)b

ABS/10 wt% DNA

Cotton/19 wt% DNA

DNA

P/N/Carbohydrates

FR materials

Polymer Substrate/ FR loading

Table 11.2 Summary of reports on bio-based FR molecules.

18 24

Yes

12 2.8 kJ/g

18.9 11.6 kJ/g

76 52 MJ/m2

Yes

Yes

Slow burning

31 36

18 28

LOI (%)g

Yes

Selfextinguishing behavior

47 31 kJ/g Yes

1.9 1.3 MJ/m2

Total heat release (THR)b

[22, 23]

[15]

[19–21]

Ref

[16]

[26]

[25]

Pass HB [24]

UL-94

280 Advanced Green Composites

PP/20 wt% PEC

Cotton/5.7 wt% CF-P/N

Cotton/8.1 wt% CF-P/N/B

Rigid PU/10 wt% Chicken feather

Chicken feather (CF)-Phosphate/ Melamine/

Borate (P/N/B)

Chicken feather

22% at 600 °C N2 2% at 600 °C Air

PET-Cotton/20 wt% Casein

7% at 900 °C Air

33% at 900 °C N2

15% at 900 °C N2

19% at 600 °C N2 4% at 600 °C Air

21% at 600 °C N2 2% at 600 °C Air

PET/20 wt% Casein

Cotton/20 wt% Hydrophobin

21% at 600 °C N2 700 °C

400–650 °C under air

CO2, CO, H2O (oxidation products)

under N2

(possible reactions) Cross-linking of aromatic rings

(futher condensations) Carbonaceous char

*R: the remaining structures in tannic acid

Figure 11.11 Postulated reaction scheme for thermal and thermo-oxidative degradation of tannic acid [96].

coating. The LOI was also significantly increased by using the combination of tannin and BA. Though applications of tannins as FR have been reported throughout the recent decades, the mechanism of thermal degradation of condensed tannins have not been carefully investigated [95, 99, 102, 103]. In contrast, tannic acid (one type of hydrolysable tannin) has a welldefined structure and has been more preferred for modeling and thermal degradation studies [96]. Fundamental understanding of thermal decomposition behavior of materials is extremely important when assessing the potential application of any material as a FR. The detailed thermal decomposition process and intumescent char-forming mechanism of tannic acid have been reported by our research group and is depicted in Figure 11.11 [96]. Tannic acid starts to decompose at about 200 °C. Through TGAFourier-transform infrared spectroscopy (FTIR) and Pyrolysis-Gas chromatography-Mass spectrometry (GC-MS) studies, 1, 2, 3-benzene triol and carbon dioxide from decarboxylation of most outer layer of galloyl units were identified as the major decomposition products both under air and nitrogen (Figure 11.12). The inner layer of the galloyl units (1, 2-benzene diol) is absent from the gas phase due to significant cross-linking reaction in the condensed phase and forms the base of the char through C-O-C linkage. The intumescence of the char was primarily formed in the temperature range of 200–400 °C due to the fast release of gas species such as carbon dioxide and phenolic monomers. At temperature lower than 400 °C, the decomposition processes of tannic acid under both air

Fire Retardants from Renewable Resources 303 CO2

nce

0.025

Absorba

0.02

OH HO

0.015

OH

0.01 0.005 40

400

0

30

300

0 Wa

200 0 um ber (c

ven

m –1 )

10

100

20 tes u in M

0

0

Figure 11.12 FTIR spectra of evolved gas from thermal decomposition of tannic acid under nitrogen [96].

and nitrogen were found to be similar indicating the oxidative atmosphere having limited influence on the primary char formation. The remaining galloyl units further cross-link and form highly intumescent graphitic char at temperatures between 400 and 800 °C. The release of non-combustible gases and formation of intumescent char upon thermal decomposition of tannic acid is promising from the perspective of utilizing tannic acid as an FR additive. The intumescent char helps protect the polymer surface from burning during combustion while the non-combustible carbon dioxide can dilute the local oxygen concentration (reducing flame propagation) and also helps with intumescent char structure. The phenolic species may also act as radical scavengers in the flame [104]. As shown in Figure 11.13, tannic acid powders exposed to propane torch for 20 seconds exhibited excellent intumescent char forming characteristics. The molten tannic acid beneath the intumescent char indicates that the char provides good thermal barrier properties. This has allowed researchers to foresee the possible protective role of tannic acid in FR coating applications. The fundamental understanding of the thermal decomposition and char forming mechanism of tannic acid evolved from this work provides better prospects for further utilization of tannic acid in FR coating and additive applications. Tannic acid has been reported to be used as char forming additive for epoxy-clay aerogel [105] and PA 6 [106]. With coating of tannic acid (2 wt%) onto epoxy aerogel foam, 20% decrease in pHRR has been achieved with a small increase of LOI value compared to neat epoxy.

304 Advanced Green Composites

Intumescent char

Molten tannic acid

Figure 11.13 Highly intumescent char from tannic acid powders after applying flame for 20 seconds.

By compounding tannic acid into PA 6, the amount of char remaining improves significantly in thermal and thermo-oxidation degradation. With 30% tannic acid loading in PA 6, the char remaining is about 13% at 750 °C in nitrogen and 16% at 550 °C in air environment with regard to 0% and 5% for pure PA 6, respectively. The intumescent char formation by tannic acid also reduces the heat release capacity (HRC) value of PA 6 by approximately 50% at 30% loading. However, the major limitation for using tannic acid as FR additive in PA is the poor thermal stability of tannic acid under the processing temperature of PA 6 (230–250 °C). Degradation of tannic acid is observed during melt-compounding. Recently we have discovered that chemical modification of tannic acid with aromatic acid chlorides through interfacial polycondensation can improve the thermal stability and char formation [107]. It was found that the chemically modified tannic acid was more thermally stable and had less discoloration in processing condition of PA 6 (weight loss 80%) of polycardanol with good thermal stability can be obtained [117]. The polycardanol obtained using this method degrades under air environment at around 250 °C and has 25% by weight char remaining at 700 °C in nitrogen environment. Polycardanol exhibits moderate flame resistance with HRC values of 300 J/g-K [118]. This synthetic strategy is easily scalable and the polycardanol obtained can be used as a char forming additive for PP. The use of a polymeric form of phenol is expected to eliminate most processing-induced decomposition problems. At 10 wt% polycardanol loading in PP, the degradation temperature of the blend increased by 36 °C as compared to neat PP resin. The char formed upon thermal degradation of the blend was found to be significant suggesting that the polycardanol may induce char formation in PP [118]. The HRC of PP blended with 15 wt% polycardanol was compared using PCFC with PP blends containing 15 wt% HBCD – a well-known commercial halogenated FR (being phased out). The HRC of blends containing polycardanol reduced by 23% and this was comparable to the HRC reduction observed for blends containing 15 wt% HBCD. The performance of this polyphenol compares favorably with some commercially available halogenated FR, purely from an HRC perspective. TGA-FTIR evolved gas analysis studies indicate that the release of nonflammable carbon dioxide and formation of char contributes to lower the HRC of polycardanol and PP-polycardanol blends. Polycardanol also seems to retard/delay the release of hydrocarbon based degradation compounds from PP upon heating. Toxicity studies carried out in accordance with the Organization for Economic Co-operation and Development (OECD) 425 guidelines indicate that Lethal Dose, 50% (LD50) values for polycardanol are greater than 2000 mg/kg [119].

Fire Retardants from Renewable Resources 307 Star-shaped cardanol oligomer (HCPP) synthesized using cardanol and hexachlorocyclotriphosphazene has also been reported [120]. HCPP undergoes thermal-initiated curing reaction. Oxygen was essential for the curing reaction, and transparent films could be easily prepared using this process. Catalysts such as cobalt naphthenate could remarkably improve the curing reaction. The cardanol-based cured films exhibited excellent thermal stability. The char yield of the cured film at 800 °C under nitrogen has been reported to be greater than 29%. These reports indicate that the fire retardancy of cardanol-based polymers can be increased by introducing phosphazene core into the structure.

11.3.4 Polydopamines Dopamine is a catecholamine neurotransmitter released by the brain that plays a number of important roles in humans and other animals. Polydopamine has been reported to exhibit strong radical scavenging activity [121] and has been used to enhance thermal stability of polymethylmethacrylate (PMMA) [104]. In FR applications, polymerized dopamine is used for PU foams [122]. The synergistic effect of char forming and radical scavenging of polydopamine in polydopamine-coated PU foam allows for imparting self-extinguishing characteristics upon exposure to flame as shown in Figure 11.15. Furthermore, this PU nanocoating foam showed 67% reduction in pHRR compared to uncoated foam. Polydopamine can also be used with APP and melamine in the synthesis of carbon foam [123]. In this case, polydopamine works as a carbon source providing intumescent char for the porous carbon foam.

Direct exposure to torch Coated foam

Control foam

Melt dripping

After removal of torch

Figure 11.15 Results of burning tests of polydopamine coated PU foam and uncoated PU foam (control) [122].

308 Advanced Green Composites

11.4

Other FR Materials from Renewable Sources

A variety of inorganic compounds such as aluminum hydroxide, magnesium oxide, calcium carbonate and antimony trioxide have been widely used as FR. These inorganic compounds slow down the propagation of flame by the endothermic decomposition process that release water or non-flammable gases and interrupting the chemical reaction sequence during burning. But these inorganic compounds are often derivatives from non-renewable mineral sources which limits their use in the future due to depletion of resources. Researchers have recently provided several solutions to renewable FR based on inorganic compounds. These are briefly discussed below.

11.4.1 Chicken Eggshell Yew et al. have recently reported their work on chicken eggshell (CES) waste as the bio-filler for intumescent FR coatings [124]. Although in this research, the flame retardancy was reported based on a system that involves three other FR additives, the addition of eggshell significantly improved fire protection performance of the coatings. CES is one of the common waste products in food industries. It consists of 95% calcium carbonate which makes it one of the least expensive renewable sources for calcium carbonate. The decomposition of CES releases CO2 that acts as a blowing agent for the formation of intumescent char. The non-flammable layer formed by calcium oxide during the decomposition protects the flammable material from fire. Yew’s work on CES as bio-filler for intumescent FR coatings provides a promising alternative to non-renewable inorganic compounds in FR system.

11.4.2 Banana Pseudostem Sap Some of the plants are known to contain abundant magnesium, aluminum, calcium, silicon and phosphorous containing mineral salts and can be utilized to impart flame retardancy as the alternative of mineral-based inorganic FR. Banana pseudostem is a waste product from banana industry and is widely available in many countries like India, Malaysia and Thailand. The liquid extract from banana pseudostem is called banana pseudostem sap (BPS) or juice. It has been reported that inorganic salts like potassium chloride, magnesium nitrate and several metal phosphates are major components in the extract of BPS. The unique composition of BPS makes it attractive to researchers. Basak et al. have reported on using BPS as an effective FR coating for various substrates including cellulosic paper, jute and cotton fabrics [125–127]. In this work, BPS treated cellulosic paper

Fire Retardants from Renewable Resources 309 achieved an LOI of 28% at a 6% weight gain [125]. An intumescent char structure was observed on BPS-treated cellulosic paper after combustion test. Compared to the ash-like char obtained on the control sample, the close-cell char structure on BPS treated paper restricted the flow of flammable gas species and prevented the propagation of flame during combustion. The flame retardancy imparted by BPS is attributed to its mineral salts. The phosphate and strong acid formed by corresponding inorganic salt during the heating catalyze the dehydration of cellulose which also accelerates the process of char formation. In the research presented by Basak et al., a fermented mixture of BPS and other nitrogen containing leaves have also been applied as bio-enriched BPS in the treatment of paper [125]. The paper treated with bio-enriched BPS achieved an LOI of 33% at an 8% weight gain. The higher amount of metallic salts present in bio-enriched BPS may act as heat sink that helps in absorbing more heat by endothermic process. Paper treated with BPS and bio-enriched BPS achieved LOI of 26% and 32%, respectively, after 300 hours of weathering under sunlight exposure [125]. Basak et al. have also expanded the application of BPS on different substrates including jute and cotton fabrics [126, 127]. The flame retardancy of BPS treated samples is summarized in Table 11.5. After the treatment of BPS in conjunction with spinach juice, the LOI and flame retardancy of cellulosic fabrics increased significantly. The vertical burning test of BPS treated jute fabrics showed excellent flame retardancy. Burning test results are shown in Figure 11.16. The dehydration and char formation process catalyzed by the metal salts and phosphate compounds are believed to contribute to the flame retardancy of treated fabrics. The durability of BPS finishing on fabric has also Table 11.5 Flammability tests of different materials coated with BPS [125–127].

Sample

Paper Paper treated Jute fabric Cotton fabric treated with biotreated treated with BPS enriched BPS with BPS with BPS

Weight gain (wt%)

6

8

8

4.5

LOI (%, treated/ control)

28/18

33/18

40/21

30/18

After flame time (treated/ control) (s)

0/40

0/40

0/60

4/60

310 Advanced Green Composites Control

Treated

Fabric in contact with flame Time = 0 sec Control

Treated

Time = 40 sec

Control

Treated

After removal of flame Time = 10 sec Control

Treated

Time = 90 sec

Figure 11.16 Comparison of burning behavior between control and BPS-treated jute fabrics at different time intervals [126].

been tested by sunlight exposure and laundry test and reported. The khaki color and flame retardancy imparted by BPS appears to be quite stable to sunlight exposure. However, after a cycle of washing, the LOI of BPS treated fabric dropped by 25% due to the good solubility of BPS and lack of strong linkage between BPS and fabric. The use of BPS or other plant extracts such as spinach juice as FR is promising because they are not only improving the sustainability of FR system but also providing a way to turn bio-based wastes into value-added products.

11.5 Prospects Since late 20th century and at the turn of the 21st century researchers have begun to explore the use of bio-derived feedstocks for the creation of safer FR mainly due to the potential toxicity and environmental persistence of some types of commonly used halogenated FR. Increased environmental consciousness as well as regulations have helped to accelerate this exploration. The combinations of naturally occurring phosphorus, nitrogen and

Fire Retardants from Renewable Resources 311 carbon-rich compounds provide possibilities for intumescent FR compositions. FR coatings prepared using sequential electrostatic LBL deposition technique have attracted a lot of attention and provided good FR efficacy. Key factors such as durability, thermal stability, and processing time (reduction in the number of bi-layers required for achieving FR characteristics) will be critical in the successful development of LBL deposition for FR coatings. Phenol-based FR exhibit promising characteristics such as char formation while potential gas phase FR action is gradually unfolding in literature reports. More detailed mechanistic study is expected to further unveil the potential FR action of polyphenols especially in gas phase. It is worth noting that the greatest challenge for bio-derived FR still stems from the complex structure and composition of biomacromolecules. Use of model compounds has been beneficial in arriving at the potential mechanism of FR action and advancing the fundamental understanding of structure-property relationships.

11.6 Summary This chapter presents an overview of some of the most promising biobased FR alternatives and the possible mechanisms of action in accordance to the current understanding. These bio-based alternatives are classified and described based on elemental composition (nitrogen and phosphorus containing compounds), chemical class (carbohydrates, nucleic acids, proteins, polyphenols) and synergistic combination of one or more of these with inorganic materials including nanoclay. Phosphorus, nitrogen and carbohydrate-based materials obtained from renewable resources, such as DNA, phytic acid, chitosan and casein, are typically utilized as intumescent FR coating. Polyphenols such as lignin and tannin exhibit good FR characteristics by forming protective char layer. To achieve comparable flame retardancy rating in the burning tests, at least 20 wt% of additives are usually required. Char formation is considered as the major mode of FR actions for polyphenol compounds. Release of non-flammable gases such as carbon dioxide can help retard the flame by dilution effect in gas phase. These bio-based additives typically facilitate in lowering the rate of flame propagation, increasing char formation and reducing the total heat released upon burning. Based on recent results from blends and coatings of polyphenols with/on polyamides and PUs, there is increasing speculation for gas phase radical scavenging action of polyphenols acting in synergy with char formation in these systems, providing self-extinguishing behavior. There are several reports of excellent FR properties being achieved at

312 Advanced Green Composites low FR loading by significantly improving the miscibility of additives in the polymer matrix. In addition to blending, sequential LBL electrostatic adsorption of polyelectrolytes from aqueous solution has been broadly explored as the ‘greener’ way of applying coating for fabrics without the use of harsh organic solvents. In several cases these renewable materials demonstrate comparable FR performance to that of commercial intumescent FR formulations. Finally, several natural extracts and even food wastes have been reported to exhibit interesting FR characteristics. Juice or sap from plants, eggshell and fruit wastes, can catalyze char formation or release non-flammable gases. Exploring the use of renewable and waste materials as feedstock for FR alternatives is quite interesting as it has the potential to unlock exciting new possibilities in low cost, safer FR materials. The abundance of natural materials that can serve as sustainable sources for the synthesis of safer FR is clearly evident.

11.7 Acknowledgements The authors would like to express their gratitude towards US Army Natick Soldier RDEC (Grant number: W911NF-11-D-0001/DO 0190/TCN 13017) for financial support. This research is approved for public release (PAO # U17–695).

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318 Advanced Green Composites 85. P. Mousavioun, W.O.S. Doherty and G. George, Thermal stability and miscibility of poly(hydroxybutyrate) and soda lignin blends. Ind. Crops Prod. 32, 656–661, 2010. 86. P. Song, Z. Cao, S. Fu, Z. Fang, Q. Wu and J. Ye, Thermal degradation and flame retardancy properties of ABS/lignin: Effects of lignin content and reactive compatibilization. Thermochim. Acta. 518, 59–65, 2011. 87. Y. Yu, S. Fu, P. Song, X. Luo, Y. Jin, F. Lu, Q. Wu and J. Ye, Functionalized lignin by grafting phosphorus-nitrogen improves the thermal stability and flame retardancy of polypropylene. Polym. Degrad. Stab. 97, 541–546, 2012. 88. A. Gregorova, B. Košíková and A. Staško, Radical scavenging capacity of lignin and its effect on processing stabilization of virgin and recycled polypropylene. J. Appl. Polym. Sci. 106, 1626–1631, 2007. 89. D. Ye, S. Li, X. Lu, X. Zhang and O.J. Rojas, Antioxidant and Thermal Stabilization of Polypropylene by Addition of Butylated Lignin at Low Loadings. ACS Sustain. Chem. Eng. 4, 5248–5257, 2016. 90. R. Zhang, X. Xiao, Q. Tai, H. Huang and Y. Hu, Modification of lignin and its application as char agent in intumescent flame-retardant poly(lactic acid). Polym. Eng. Sci. 52, 2620–2626, 2012. 91. B. Prieur, M. Meub, M. Wittemann, R. Klein, S. Bellayer, G. Fontaine and S. Bourbigot, Phosphorylation of lignin to flame retard acrylonitrile butadiene styrene (ABS). Polym. Degrad. Stab. 127, 32–43, 2016. 92. L. Liu, G. Huang, P. Song, Y. Yu and S. Fu, Converting Industrial Alkali Lignin to Biobased Functional Additives for Improving Fire Behavior and Smoke Suppression of Polybutylene Succinate. ACS Sustain. Chem. Eng. 4, 4732–4742, 2016. 93. H. Zhu, Z. Peng, Y. Chen, G. Li, L. Wang, Y. Tang, R. Pang, Z.U.H. Khan and P. Wan, Preparation and characterization of flame retardant polyurethane foams containing phosphorus–nitrogen-functionalized lignin. RSC Adv. 4, 55271–55279, 2014. 94. A. Arbenz and L. Avérous, Tannins: A Resource to Elaborate Aromatic and Biobased Polymers, in: Biodegradable and Biobased Polymers for Environmental and Biomedical Applications, S. Kalia and L. Avérous (Ed.), pp. 97–148, John Wiley & Sons, Inc., 2016. 95. H. Tributsch and S. Fiechter, The material strategy of fire-resistant tree barks. WIT Trans. Built Env. 97, 43–52, 2008. 96. Z. Xia, A. Singh, W. Kiratitanavit, R. Mosurkal, J. Kumar and R. Nagarajan, Unraveling the mechanism of thermal and thermo-oxidative degradation of tannic acid. Thermochim. Acta. 605, 77–85, 2015. 97. K. Khanbabaee and T. van Ree, Tannins: Classification and definition. Nat. Prod. Rep. 18, 641–649, 2001. 98. G. Tondi, A. Pizzi and R. Olives, Natural tannin-based rigid foams as insulation for doors and wall panels. Maderas Cienc. Tecnol. 10, 219–227, 2008. 99. M.A. Pantoja-Castro and H. González-Rodríguez, Study by infrared spectroscopy and thermogravimetric analysis of Tannins and Tannic acid. Rev. Latinoam. Quím. 39, 107–112, 2011.

Fire Retardants from Renewable Resources 319 100. G. Tondi, S. Wieland, T. Wimmer, M.F. Thevenon, A. Pizzi and A. Petutschnigg, Tannin-boron preservatives for wood buildings: Mechanical and fire properties. Eur. J. Wood Wood Prod. 70, 689–696, 2012. 101. G. Tondi, L. Haurie, S. Wieland, A. Petutschnigg, A. Lacasta and J. Monton, Comparison of disodium octaborate tetrahydrate-based and tannin-boronbased formulations as fire retardant for wood structures. Fire Mater. 38, 381–390, 2014. 102. M. Gaugler and W.J. Grigsby, Thermal degradation of condensed tannins from Radiata pine bark. J. Wood Chem. Technol. 29, 305–321, 2009. 103. S. Nikkeshi, Agent for imparting flame retardancy to thermoplastic resin, US Patent 6624258, assigned to Tohoku Munekata, 2003. 104. J.H. Cho, K. Shanmuganathan and C.J. Ellison, Bioinspired catecholic copolymers for antifouling surface coatings. ACS Appl. Mater. Interfaces. 5, 3794– 3802, 2013. 105. X. Lang, K. Shang, Y.-Z. Wang and D.A. Schiraldi, Low flammability foamlike materials based on epoxy, tannic acid, and sodium montmorillonite clay. Green Mater. 3, 43–51, 2015. 106. W. Kiratitanavit, Z. Xia, A. Singh, R. Mosurkal and R. Nagarajan, Tannic Acid: A Bio-based Intumescent Char-forming Additive for Nylon 6, Society of Plastics Engineers. in: ANTEC 2016, P046, Indianapolis, 2016. 107. Z. Xia, W. Kiratitanavit, R. Mosurkal, J. Kumar and R. Nagarajan, Chemically modified tannic acid: A bioderived char forming additive for nylon 6 (GCE4) – ACS Presentations on Demand. 19th Annual Green Chemistry & Engineering Conference. Washington D.C., 2015. 108. Z. Xia, W. Kiratitanavit, S. Yu, P. Facendola, S. Thota, J. Kumar, R. Mosurkal and R. Nagarajan, Modified tannic acid – A bioinspired fire resistant char forming additive for polyamide. Presented at the 252nd ACS National Meeting, Philadelphia, Pennsylvania August 21, 2016. 109. S. Nam, B.D. Condon, M.B. Foston and S. Chang, Enhanced thermal and combustion resistance of cotton linked to natural inorganic salt components. Cellulose 21, 791–802, 2014. 110. S., Nam, H.J. Kim, B.D. Condon, D.J. Hinchliffe, S. Chang, J.C. McCarty and C.A. Madison, High resistance to thermal decomposition in brown cotton is linked to tannins and sodium content. Cellulose. 23, 1137–1152, 2016. 111. S. Nam, B.D. Condon, Y. Liu and Q. He, Natural resistance of raw cotton fiber to heat evidenced by the suppressed depolymerization of cellulose. Polym. Degrad. Stab. 138, 133–141, 2017. 112. S. Nam, B.D. Condon, Z. Xia, R. Nagarajan, D.J. Hinchliffe and C.A. Madison, Intumescent flame-retardant cotton produced by tannic acid and sodium hydroxide. J. Anal. Appl. Pyrolysis. 126, 239–246, 2017. 113. H. Ejima, J.J. Richardson, K. Liang, J.P. Best, M.P. van Koeverden, G.K. Such, J. Cui and F. Caruso, One-step assembly of coordination complexes for versatile film and particle engineering. Science 341, 154–157, 2013. 114. J. Guo, Y. Ping, H. Ejima, K. Alt, M. Meissner, J.J. Richardson, Y. Yan, K. Peter, D. von Elverfeldt, C.E. Hagemeyer and F. Caruso, Engineering

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12 Green Composites with Excellent Barrier Properties Arvind Gupta, Akhilesh Kumar Pal, Rahul Patwa, Prodyut Dhar and Vimal Katiyar* Department of Chemical Engineering, Indian Institute of Technology Guwahati, Guwahati Assam, India

Abstract This chapter discusses the recent technological advancements and innovations in the development of novel biodegradable polymers and nanofillers derived from renewable feedstock with a strategy to convert ‘waste into wealth’ with special focus on green composites for stringent food packaging. The biopolymers have been classified into various categories depending upon the source of origin and their physicochemical and structural properties are discussed. A wide pool of bionanofillers derived from biomass has been reported with their strategic modifications through various routes in developing high gas barrier films. With increased environmental concerns and enforcement of strict legislations, replacement of conventional polymers with biodegradable ones has become need of the hour. However, with some of the petroleum-based polymers with biodegradable character, may be effectively commercialized only when industrially viable processes are developed along with its wide scale adaptations in the society. Therefore, this chapter presents appropriate use of bio-derived bionanofillers and their green composites and packages using selective polymers, derived from both from bio-based and petroleum feedstock; for targeted high gas barrier applications especially in areas related to packaging. Keywords: Biodegradable polymers, biopolymers, nanocomposites, packaging

12.1 Introduction According to the report published by Ellen MacArthur foundation at the World Economic Forum, 2016, there will be more accumulated waste *Corresponding author: [email protected] Anil Netravali (ed.) Advanced Green Composites, (321–368) © 2018 Scrivener Publishing LLC

321

322 Advanced Green Composites plastic in the sea than marine creatures by 2050 [1]. Faced with such realities, increased environmental concerns and stricter rules and regulations have been enforced by legislations around the world towards the use of conventional plastics. Fortunately, with growing research and technological advancements towards the development of biodegradable polymers, it is now possible to selectively replace petroleum-based commodity plastics with the newly developed biodegradable ones. In addition, drastic depletion in fossil-fuel reserves and hike in price of petroleum based plastics in the past few years has led the commercialization of biopolymers at an accelerated pace via techno-economically feasible approach. Another disadvantage of the petroleum-based polymers is that they take long time, typically hundreds of years, to degrade which leads to their accumulation in the ecosystem. Alternatively, their incineration produces large volumes of greenhouse as well as toxic gases [2]. On the other hand, bioplastics undergo easy degradation into non-toxic fragments after their intended use [3]. The bioplastics are usually classified into two types, firstly, ‘biopolymers’ derived from microbes which are produced through fermentation process utilizing the waste agricultural-residues and secondly, ‘biobased polymers’ which must have at least ~30% of its carbon backbone derived from renewable resources synthesized through polymerization process [4]. However, those bioplastics are specifically termed as ‘biodegradable polymers’, where ~90% of their carbon backbone disintegrates into non-toxic monomeric units (such as carbon dioxide, nitrogen or ammonia, water, inorganic/organic derivatives, etc.) in soil or composting conditions within a stipulated timeline of ~180 days. There are different governing bodies such as ASTM standard D-5488-94d and European norm EN13432-2000, which regulate the extent of biodegradability of polymers under aerobic or anaerobic conditions. The various degradation routes of bioplastics can be composting, enzymatic degradation and/or biotic/microbial degradation [5–7]. The standards also take into consideration the degree of mineralization as well as the type of residues generated during biodegradation process, especially evaluating its eco-toxicity potential. Biodegradation of the polymers at such enhanced rates prevents its accumulation in the environment. Also, the degraded monomeric units are mostly benign biological molecules which are used up or metabolized by microbes and plants in their normal biogeochemical cycles. The monomeric fragmented units formed during decomposition processes are mostly carbon-rich materials especially organic units containing short oligomeric chains or humus-like substances and gaseous carbondioxide or ammonia. Such small molecules can undergo dissolution in soil conditions enriching it for easy uptake by photo-synthetically active

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organisms as food source in their metabolic cycle. In their nutrient cycle, these degraded components are converted into value-added products making the life cycle assessment of the biopolymers an environmentally benign and economically viable approach [8, 9]. As a result, through these biodegradation routes, bioplastics are able to enter into the carbon cycle, by reducing both CO2 generation as well as polymeric waste accumulation. The occurrence of biodegradation is completely absent in synthetic non-degradable plastics which are almost inert to external weathering conditions as well as enzymatic action from microbes. However, recent reports on degradation of poly(ethylene terephthalate) (PET) into small fragments by newer bacterial strains seems quite promising [10]. There are still some techno-economical issues related to biopolymers such as timeframe for degradation, economic viability and scalability which needs to be addressed. In this respect, biopolymers can help to overcome such challenges and also have the potential to revolutionize the packaging industry with their inherent advantages over petroleum based synthetic plastics such as biodegradability, sustainability, non-toxicity and abundance of feedstock.

12.2

Biodegradable Polymers: Classifications and Challenges

In recent days, biodegradable polymers have become the most studied class in the field of polymer science, which have comparable properties with synthetic polymers. Depending on the source of origin, biodegradable polymers can be classified as either derived naturally through fermentation or by extracted biomass based precursor through polymerization [11–13]. There are some classes of aliphatic/aromatic co-polyesters derived from petroleum based resources which are degradable under microbial action as well as composting conditions and, hence, are categorized within the biodegradable polymers. However, some polymers such as bio-PE (polyethylene) and bio-PET (polyethylene terephthalate) have been recently developed from bio-based sources but are non-biodegradable which is still a major concern [14, 15]. Various categories of biodegradable polymers derived from agricultural sources such as polysaccharides or protein based polymers and biopolyesters which are usually extracted from micro-organisms through biotechnological routes or derived from petroleum derivatives as shown in Figure 12.1. The first category of biopolymers are those which are derived directly from bio-sources and consist of polysaccharides, lipids and proteins. Polysaccharides that include cellulose, chitosan, gum,

324 Advanced Green Composites Biodegradable polymers Directly extracted from biomass

Produced by microbes

Polymerization routes

Polysaccharides PHB Cellulose

Chitosan

Starch

PLA

Bacterial cellulose

Lipids Xanthan Proteins

Figure 12.1 Biodegradable polymers with various source of origin derived from different renewable resources.

starch, pectin, etc., are the most abundantly available and naturally derived biopolymers. These are generally extracted from renewable lignocellulosic biomass, sea creatures and agricultural resources [16, 17]. Collagen, silk, keratin and gelatin belong to the class of proteins which are generally isolated from animal extracts, insects, worms, animal hair, bird feathers, etc., through chemical treatment [18]. Recently, research on bio-derived polymers have been widely carried out to develop strategies to convert ‘wastes into wealth’ due to their interesting physico-chemical properties such as higher structural properties, bio-origin, high functionality, non-toxicity and improved bio-compatibility. Extraction of renewable biopolymers followed by its surface modifications and successive development of nanoparticles have found potential applications in theranostics, electronics, drug delivery, implants, packaging and energy storage devices. In this chapter we have discussed various aspects of such biopolymers with a focus on their nano-derivatives and their strategic chemical modifications to develop high barrier green composites for targeted applications in packaging. The second category of biodegradable polymers such as polylactic acid (PLA) and polyhydroxyalkanoates (PHAs) are generally derived through polymerization or microbial fermentations. PLA is one of the most studied bio-derived polymer which can be synthesized via condensation polymerization of lactic acid which itself is derived through microbial fermentation of biomass. PHAs are a class of biopolymers which are directly accumulated by the microbes as a response to certain physiological stress conditions. Hydroxyalkanoates exist in eleven derivatives of its monomeric units, which assemble to form

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homo, co-and block-copolymers depending upon loading and fraction of the monomers [19]. The physico-chemical properties of such co/heteropolymeric structures of PHAs, depend upon the type of bacterial strains, synthesis route followed and variation of the carbon/nitrogen ratio in the nutrient media. Comparison of barrier properties and structural behavior of PLA and PHAs with the petroleum derived polymers such as PP, PET etc. has been shown in Table 12.1. It is observed that pristine biopolymers have poor gas barrier and thermo-mechanical properties in comparison to conventional petroleum polymers thus hindering their possible usage towards packaging applications. It is noteworthy to mention that naturally occurring polymers such as cellulose, chitosan and silk have superior mechanical properties but show poor film forming ability and due to their hydrophilic nature they also show lower stability under ambient conditions. Moreover, processing of the biodegradable polymers through the industrially viable processes such as extrusion or blow molding have proven to be challenging due to drastic reduction in their molecular weight. In order to improve the properties of PLA, poly (butylene succinate) (PBS) and PHAs, several varieties of nanofillers such as metallic, inorganic as well as organic nanoparticles such as polymeric, carbon-based, clay or ceramic based nanomaterials have been used [20]. Incorporation of such broad range of nanomaterials depending on the targeted applications have led to significant improvements in the structural, gas barrier and thermo-mechanical behavior of biodegradable polymers [21]. However, there remain several challenges on utilizing the above mentioned polymeric nanocomposite films as packages for storage of food products due to their biodegradability, cytotoxicity issues and unknown migration behavior when in contact with food items. Bionanoparticles have the potential to be used as reinforcing agent in the fabrication of green bionanocomposites, as a way to overcome such limitations. Introduction of such biofillers retains the integrity of polymers, especially their non-toxicity and biodegradable nature. This has led to the development of “Green and Sustainable” bionanocomposites. However, commercialization of such biodegradable polymers as viable products has been limited due to high production cost, low yield as well as their nonrecyclability unlike petroleum-derived polymers. Therefore, researchers have focused on the development of industrially viable processing techniques for such biodegradable polymers with improved properties as well as easy recyclability which reduces the product costs. In subsequent section, we discuss various types of biopolymers and their derivatives used as bio-nanofillers in the production of green and biodegradable nanocomposites with improved physico-chemical and gas barrier properties required for packaging application.

Source Biopolymers Cocoon silk (B.mori) Spider dragline silk Spider Viscid silk (fiber) Chitosan film Silk film CNCs or CNFs Carboxylate CNFs Petroleum derived polymers High Tenacity Rayon High Tenacity Nylon Polyethylene High-tensile steel Polypropylene Kevlar 49TM

900 22-31 1500 336 3600

600 1100 44 67.7 ~200

Tensile strength (MPa)

80 6 50

-

-

-

0.8 0.1

70 160 -

Toughness (MJ·m−3)

0.6 0.7-2.3 0.1-0.5 -

Tenacity (GPa)

12 20 1200 0.8 345 2.7

14 ≤30 ≤300 21 5.5 -

Elongation (%)

13 4 1.0 1.8 -

6 9.5-30 ≤1 324 1.9 150

Modulus (GPa)

76 cc.μm/m2.day.atm 50-200cc.μm/m2.day.kPa 3900 cc.μm/m2.day.atm -

0.6cc.μm/m2.day.kPa 0.0006cc.μm/m2.day.kPa

OTR

[29] [23] [30] [31] [32] [33]

[22] [23] [24] [25] [26] [27] [28]

Ref

Table 12.1 Mechanical and gas barrier properties of various biopolymers and biodegradable polymers along with its comparison with petroleum derived polymers.

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Biodegradable polymers [34] PLA 28-50 2-6 1.2-3 10-30 cc.μm/m2.day.kPa [19] PHA 17-104 2-1000 0.3-0.4×10-16m3.m/m2.s.Pa PBS 25 175 0.268 0.2×10-10 cm3cmcm-2s-1Pa-1 [35] [36] [37] PTT 50 160 2.09 2.9cm3. cm/m2·day.atm Stereo-PLA (Solution cast film) 49.0 --10 1.8 -[38] Stereo-PLA (melt spun) 335 ---4.7 -[39] Stereo-PLA (solution cast) 410 --21 9.5 -[40] Bio-based and biodegradable nanocomposites [41] PLA/PBS blends 104.5 228 3.5 2×10-15cm3cmcm-2s-1Pa-1 PLA/modified Chitosan films 40.3 34.3 0.490 1.83 cc.mm/m2·day·kPa [25] [42] PLA/CNC films 65 5 1.3 3.9×10-2 cm3.mm.m-2.day.kPa PLA/Silk films 33 3 3.4 [43] PHB/CNC films 57 6 1.41 0.1×10-16 m3.m/m2.s.Pa [44] [45] PLA/TMC-328 (melt extruded) -----1.99×10-20 m3.m.m-2.s-1.Pa-1 -15 3 -2 -1 -1 [46] Stereo-PLA/Graphene oxide -----1.5×10  cm .cm.cm .s .Pa (Solution cast) CNCs: Cellulose Nanocrystals, CNFs: Cellulose nanofibers, PTT: Poly(trimethylene terephthalate), PLA: Poly (Lactic acid),PHA: Polyhydroxyalkanoates, PHB: polyhydroxybutyrate, PBS: Polybutylene succinate, TMC-328: N,N ,N -Tricyclohexyl-1,3,5-benzene-tricarboxylamide.

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328 Advanced Green Composites

12.2.1 Poly (lactic acid): Properties Evaluation, Modifications and its Applications Among all the biobased polymers, PLA is the most promising candidate which can replace some of the conventional polymers [47]. The biggest advantage of PLA is that, after its service life, it can microbially degrade into carbon dioxide, methane and water over the period of several months to few years [48], whereas, most petroleum based polymers take hundreds of years to degrade and go back to the nature. PLA is a bio-based, biodegradable, bioabsorbable polymer synthesized by melt polycondensation of lactic acid derived oligomers. In eighteenth century, Swedish chemist Carl Wilhelm Scheele reported the isolation of lactic acid from the sour milk [49]. Lactic acid is a chiral molecule which exists in the form of L and D isomers and produced by the fermentation of sugar. PLA may be produced using several polymerization techniques such as condensation polymerization (polycondensation), ring opening polymerization (ROP) and azeotropic dehydrative condensation polymerization. The polycondensation and ring opening polymerizations are the techniques which are mainly used for the production of high molecular weight PLA. Using polycondensation, high molecular weight PLA can be produced after solid state polymerization. However, it has limitation of time and energy consumption. High molecular weight PLA with higher conversion in relatively less time can be synthesized by ROP of lactide which is dimmer of lactic acid. The process for lactide synthesis was first patented in 1913 by Reinhold Gruter [50] and in 1932, ROP of lactide was reported by Wallace H. Carothers [51]. Earlier, due to instability in humid condition, PLA was not considered as a useful material. In 1966, PLA found its place in biomedical application due to its biodegradability, bio-absorbability, biocompatibility and non-toxic nature [52]. In last two to three decades, the use of PLA in biomedical application has increased due to the earlier mentioned properties and subsequently, it has also found its place in high strength applications such as textiles, agriculture, electronics, packaging and other applications. PLA also has comparable thermal, mechanical and gas barrier properties which makes PLA more promising candidate to replace petroleum based polymers. However, PLA also has drawbacks such as relatively low glass transition temperature, low heat deflection temperature, lower melting temperature, slow crystallization, lower melt elasticity and relatively poor barrier properties. While improvements in these properties have been challenging for the scientific community, they can be tuned by development of composites with designed green polymer matrices. Several methods such as developing biocomposites, polymer blending or formation of stereocomplexes (SC) have

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been adopted by the researchers to improve the structural and physicochemical properties of PLA so that it can compete with the properties of petroleum derived polymers. Development of PLA based biocomposites provides an excellent opportunity to fabricate novel biodegradable polymeric films which can effectively improve its intrinsic properties with good potential for commercialization. It has been proven that mixing of PLA with fillers or polymers such as cellulose [53], clays [54], graphene [55], chitosan [25] gums [56], cellulose nanocrystals [57, 42], silk [58] etc. improves its properties and also yields ecological as well as economical composites. Enhanced gas and moisture barrier properties of PLA are very crucial in packaging application in order to improve the shelf life of the food product. Various methods including layer by layer (LbL), surface treatment, blending, etc., can be used for the fabrication of biocomposites. Depositing desired number of multilayers of biopolymer in combination with other low oxygen solubilizing additives may reduce the oxygen permeability (OP). Fukuzumi et al. [59] prepared the 2,2,6,6-tetramethylpiperidine1-oxyl radical (TEMPO)-oxidized cellulose nanofibers (TOCN) and coated the PLA substrate with them. Depositing fifty layers of TOCN on the surface of PLA significantly reduced the OP to about 1/750 times. Aulin and his group [60] used LbL method to prepare multilayer of polyethylenimine (PEI) and nanofibrillated cellulose (NFC) or carboxymethyl cellulose (CMC) on the surface of PLA substrate. They found that the PLA films coated with 50 bilayers of PEI/NFC and PEI/CMC exhibited OP (~10 cm3/m2.day.atm) were significantly lower (by more than an order of magnitude) than uncoated PLA (151.5  cm3/m2.day.atm) at 23 °C and 50% RH. In addition, the coated PLA films became less sensitive to moisture after LbL deposition method. They also found that the film was more than 80% transparent to visible light after coating. The LbL method was also employed by Laufer et al. with the objective to reduce the OP of PLA which could slow down the oxidative degradation and ultimately improve the shelf life of the food [61]. They developed the method in which the PLA film needed to be sequentially dip coated with chitosan and montmorillonite (MMT) clay solution followed by drying. After coating 30 times with chitosan and clay, the OP of PLA was reduced from 177.2 × 10–16 cm3. cm.cm–2.s–1.Pa–1 to less than 0.03 × 10–16  cm3.cm.cm–2.s–1.Pa-1, more than 3 orders of magnitude. They postulated that the layer of clay acted like a nanobrick wall film which built extremely tortuous path for the oxygen molecules. However, the commercialization of such processes can be debated as these may have significantly higher processing costs. Since the space between clay layers improves the barrier property significantly [62],

330 Advanced Green Composites the researchers used chitosan as spacer between clay layers. Chitosan itself believed to have low barrier properties but its permeability is far higher than PET (17.3 × 10–16 cm3.cm.cm–2.s–1.Pa–1). They concluded that addition of bare clay or chitosan layer to PLA film do only moderate reduction in O2 permeability. They also noted that conventional thick film composites have limitation of filler loading, close to 10 wt%, but use of LbL process can create composites with 90 wt% clay without affecting their transparency and mechanical properties. The gas or water vapor barrier properties of polymers can also be improved by surface treatment processes such as UV/ozone treatment [63], temperature treatment [64], plasma radiation [65], hyperthermal hydrogeninduced cross-linking (HHIC) technology, etc. The UV/ozone treatment is also used for the water treatment, food sterilization, etc. Due to the ozone treatment, the polymeric surfaces become more hydrophilic. The temperature or heat treatment method can increase the crystallinity of the polymer due to the formation of spherulites which reduces the diffusivity. As in the case of nanoclay, these crystals increase the tortuous path for the diffusing gas molecules which ultimately reduces the gas permeability, a characteristic highly desirable for stringent food packaging applications. Plasma treatments can be an effective method to modify the polymeric surfaces in order to get improved gas barrier properties of plastics and their composites. Plasma treatments of surfaces using fluorinated compounds and sulfur hexafluoride can result in increased hydrophobicity of the surface. The fluorinated functional groups are apolar in nature and surfaces such as PLA, modified by these functional groups, experience reduced surface energy which leads to increased contact angle with water. As a result, the first step during the permeation process can be a limiting step that may slowdown the interfacial solubility of molecular water. In the food industry, to have long shelf life and maintain the quality of foods, high gas barrier property of packaging material is an important parameter to protect the food from moisture. It has been found that some plasma treatments improve the water repellent properties but cannot prevent the bulk transport of moisture across the film [5]. For such situations multilayer film preparation could be the alternative solution [65]. HHIC technology has been used for selective breaking of C-H or Si-O bonds, in the absence of any chemical reagents, without affecting other atomic bonds. In this method carbon radicals are generated due to breakage of C-H bonds which promote the cross-linking (Figure 12.2). The cross-linking induced by hyperthermal hydrogen is driven by plasma. This plasma is generated using hydrogen gas and protons are extracted and electric field is used to give desired acceleration. The high energy protons collide with molecular

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Cleavage of C-H bond by collision HHIC reactor

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Figure 12.2 Schematic of surface cross-linking of the substrate induced by hyperthermal hydrogen collision (HHC) technique.

hydrogen and produce the hyperthermal molecular hydrogen which carries no electric charge. The bombardment of the hyperthermal molecular hydrogen on the substrate cleave the C-H or Si-H bonds due to the kinematic selectivity of energy deposition on the hydrogen atom present on the substrate. The breakage of each C-H bond induces the cross-linking reaction which produces a stable controlled cross-linking surface or molecular layer on the substrate. Du et al. used the HHIC technology to cross-link the PLA film surface in order to reduce its water vapor permeability (WVP) [66]. Due to hyperthermal hydrogen collision, the C-H bonds of the PLA molecules were selectively cleaved without breaking other bonds. This resulted in cross-linking of carbon radicals generated from the organic molecules. The cross-linked layer served as a barrier and reduced the WVP of PLA from 26.7 to 22.9 gm/m2.day and improved its hydrophobicity. This technology can also be used for the development of polymer laminates [6]. Generally, clays have layered structure which results in a tortuous path for gas molecules even at low loading, thus reducing the gas diffusivity through the matrix. A review containing detailed information on polymer clay composites for gas barrier properties has been provided by Cui et al. [67]. Blending of polymers is another approach by which the gas barrier properties can be significantly improved. Armentano et al. found that addition of 15% PHB into PLA matrix increases its crystallinity due to enhanced crystal packing thereby reducing the OP by about 35% [68]. With the objective to reduce the PLA’s gas barrier properties, Kakroodi et al. developed a technique to form in situ microfibrils of polyamide (PA)

332 Advanced Green Composites on the surface of PLA by melt blending followed by hot stretching at the die exit, using a take-up roller [69]. They found that the addition of just 3 wt% of PA led to significant reduction of about 38% in the OP. OP reduction was found to be directly related to the crystal size and its distribution. In this case, PA microfibrils acted as nucleating agent enhancing the nucleation density leading to the production of large amount of small crystals. Ultimately, these small crystals constrain the chain mobility in the amorphous region and tight packing of chains leads to the reduction in OTR. Recently, new type of crystals called ‘SC crystallites’ have been formed, when enantiomeric PLA such as poly(L-lactic acid) (PLLA) and poly-(D-lactic acid) (PDLA) are mixed together. The melting temperature of SC PLA (sPLA) was found to be 50 °C higher than enantiomeric pure PLA [70]. A detailed schematic about the production and application of PLLA, PDLA and sPLA is presented in Figure 12.3. Stereocomplexation is a special type of crystal formed by interaction between two enantiomeric PLA chains which is responsible for the improvement in the thermal, mechanical and barrier properties of the end products [7-10, 64-66]. Enantiomeric pure PLA adopts 103 helical configuration whereas sPLA having 31 helical configuration responsible for densely packed polymer chains which improves the various properties of PLA [71]. Xu et al. fabricated the graphene oxide (GO) based flexible biopolymer films [72]. They found that during solution mixing, SC nanospheres were forming and surrounding the GO sheets which acted as nucleating sites for spherulite development. Presence of GO in the matrix drives the steady increase in SC concentration and suppressed the formation of homocrystals. The OP of the composite films was reduced from 4.869×10-15 cm3.cm.cm-2.s-1.Pa-1 to 0.2553×10-15 cm3.cm.cm-2.s-1.Pa-1 by the addition of 0.5 wt% GO and further reduced to 0.06264×10-15 cm3.cm.cm-2.s-1.Pa-1 after annealing. Presence of GO was seen to form impermeable nano-barrier walls which improved the resistance to diffusing gas molecules. In addition, enhancement in the SC crystallinity in the matrix also played a significant role in the reduction of oxygen solubility. The production of PLA grafted fillers by in situ polymerization is another technique to improve the dispersion of fillers in polymeric matrix. In this method, the PLA grafted fillers are synthesized by polymerization of L-lactide in the presence of fillers and subsequently, filler grafted PLA are dispersed into pristine PLA matrix. The grafting of PLA chains on the surface of filler influences the interaction between the matrix and filler which leads to the homogeneous dispersion. Re et al. demonstrated that a dual approach, i.e. use of filler and stereocomplexation, significantly overcomes

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Figure 12.3 Different routes for production of PLA and its stereocomplexation for potential applications in various sectors.

the drawbacks related to properties of polymeric matrix [73]. They first prepared and purified the PLLA grafted clay (Cloisite 30B) and then melt blended it with PDLA to get the homogeneous dispersion of grafted clay followed by the formation of SC leading to 46% reduction in the OP of the film. From the above discussions it can be concluded that the structural and barrier properties of PLA may be enhanced significantly, through various chemical modification approaches to make them almost comparable to those of conventional petroleum-derived polymers.

12.2.2 Cellulose Based Composites: Chemical Modifications, Property Evaluation, and Applications Cellulose is the most abundantly available natural biopolymer, commonly extracted from plant-based renewable resources. Due to the high chemical functionality, its backbone can be modified through various routes depending upon the target applications. Cellulose products are composed of both amorphous and crystalline domains in varying ratios depending upon the source of origin. The crystalline domains can be extracted through controlled acid hydrolysis by selectively removing the amorphous segments and the product isolated is termed as cellulose nanocrystals (CNCs) [74]. The morphology and dimensions of such crystalline domains vary with the type of biomass selected, cellulose polymorph, type of pretreatment and hydrolysis process followed for its extraction, which ultimately contributes towards the mechanical properties [75]. CNCs usually have rod-like

334 Advanced Green Composites morphology with tunable aspect ratio along with high hydroxyl functionality making them an ideal reinforcing agent when dispersed in biodegradable polymers such as poly(lactic acid) (PLA) [42], polyhydroxybutytrate (PHB) [44], polycaprolactone (PCL) [76], PBS [35], etc. Along with the aforementioned properties, CNCs are non-toxic, biocompatible, possess high surface area and self-ordering behavior which leads to some unique properties. They consist of parallely stacked crystalline segments which significantly improves anisotropy leading to improvements in mechanical behavior with high elastic modulus of |between 110 and 220 GPa along the axial direction and between 2 and 50 GPa along the traverse direction [77]. The aforementioned mechanical properties of the CNCs are known to significantly depend upon the source of biomass origin, pretreatment method followed by the extraction of cellulose, type of acid hydrolysis methodology selected, cellulose polymorphism and the functionality. Presence of such inherent characteristics in CNCs can be expected to improve the structural behavior of prepared nanocomposites when dispersed under ideal conditions. Moreover, alignment of the CNCs through application of external forces in polymeric matrix leads to anisotropic mechanical, thermal, electrical and optical properties that are tunable [78, 79]. CNCs have also been used as coating materials, through LbL approach, for various polymeric substrates using different processing techniques such as dip coating, spin coating, spraying or electrochemical depositions for improving the barrier properties [80, 28, 60]. CNCs are known to form a percolated network like structure when dispersed in polymeric matrix, which thereby provides a tortuous pathway to the permeation of oxygen and water vapor molecules across the nanocomposites [81]. The hydroxyl groups present in CNCs are known to form inter-molecular hydrogen bonding with the polymeric matrix, which further restricts the solubility of the permeate gases in the fabricated nanocomposite films. Such films, developed with improved thermo-mechanical and barrier properties using CNCs as nanofillers, can be potentially used for food packaging applications due to its non-toxicity and migration values within the prescribed limits as per the standard legislations [19]. Therefore, CNCs as nanofillers can be expected to synergistically reduce the diffusivity as well as solubility of the gases leading to significantly improved gas barrier along with enhanced structural properties of the bionanocomposites. As discussed in the previous section, CNCs have potential to be used as nanofillers due to their unique morphology and abundant functional groups on their surface for improving the barrier properties of biodegradable polymers [42]. However, several challenges remain in obtaining enhanced dispersion of the hydrophilic CNCs which tend to agglomerate

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Permeation of O2(g)/H2O(v)

Permeation of O2(g)/H2O(v)

Permeation of O2(g)/H2O(v)

in hydrophobic polymeric matrix materials. Presence of such agglomerated CNCs in the fabricated nanocomposite films prevents the effective transfer of desired characteristics of CNCs to the polymeric matrix materials. CNCs can be added to polymers either directly or after chemical modifications of their surfaces. CNC surface can be modified through grafting by radical polymerization or by reactive processing. The permeation of gas molecules across the polymer/CNC nanocomposites is expected to be dependent on several factors, namely: (i) morphological dimensions such as length/width and aspect ratio, (ii) degree of crystallinity and (iii) interfacial interactions between the CNCs and the polymeric matrix. This is shown in the schematic presented in Figure 12.4. The transport of molecules through the polymeric CNC nanocomposite membranes comprises of following three distinct steps; adsorption of molecule onto the surface of film (first step), followed by its diffusion through the membrane (second step) and finally desorption from the surface of the films (third step) [82]. In the first step, the permeate moieties

CNCs

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Cellulose nanocrystals (CNCs) Magnetic cellulose nanocrystals (MGCNCs) Cellulose nanofibers (CNFs) Polymeric matrix

Figure 12.4 Effect of different varieties of nanocelluloses on the barrier properties of films.

336 Advanced Green Composites should first dissolve onto the film boundary surface, rate of which depends upon the adhesion between the moiety and the functional groups present on the polymer nanocomposite films/membranes. Subsequently, in the second step, the rate of diffusion of permeate moieties will depend on the tortuosity (of the path) effect as well as interfacial surface interactions between the CNCs and the polymer. The tortuosity effect attributed by CNCs, will depend upon their morphology, aspect ratio, degree of inter/intra-molecular hydrogen bonding and its ability to form dense percolated network with the polymeric matrix. The nanocellulose usually exists in different morphologies such as CNCs or nanowhiskers (CNWs) with rod-like morphology, cellulose nanofibrils (CNFs) with fiber-like morphology or cellulose nanoparticles with spherical morphology (as shown in Figure  12.4), which have wide variation in crystallinity ranging from 40–90% [83]. CNFs, with fiber-like morphology have relatively lower crystallinity but have the ability to form dense network-like structure through entanglements and formation of the inter/intra-molecular hydrogen bonding [84]. Presence of such stable network of fibers nullifies the effect of its low crystallinity and provides high resistance for gas permeation compared to rod-like CNCs. Poor adhesion leads to formation of voids/pores at the interface of CNCs and polymer matrix, which provide pathways for easy permeation of gas molecules. Even though CNCs or CNFs provide high barrier towards permeation of gas molecules, their effect towards water vapor permeation doesn’t have such significant effect [85]. This is probably due to the hydrophilic nature of CNCs and CNFs which results in adsorption of water vapor molecules under humid conditions leading to swelling of the surface. However, such problems can be overcome through strategic surface modifications of CNCs (as discussed in subsequent section), which reduces the water adsorption capability of the crystalline domains in cellulose, thereby, decreasing the water vapor transmission rate across the films. CNF based films are generally produced through vacuum filtration followed by air drying under compression. The films prepared from CNFs, obtained through TEMPO oxidation of cotton, are usually transparent with crystallinity in the range of 75–90%, which leads to significantly high oxygen barrier properties (~17±1 ml.m-2day-1) which are interestingly comparable with the petroleum derived polymers such as ethylene vinyl alcohol and polyvinylidene chloride. Surface modified bacterial cellulose membranes through esterification process leads to significant reduction in the permeation of oxygen, carbon dioxide and nitrogen gases by as much as 50% [86]. However, several challenges remain in fabrication of

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the nanocellulose based films as packaging materials due to their highly brittle nature as well as high water adsorption capacity which almost leads to disruption of its structural stability. High barrier membranes made up of cross-linked PVA/CNC nanocomposites with poly(acrylic acid) (PAA) through simple heat treatment provides enhanced resistance to water vapor and trichloroethylene permeation. It was observed that at 10 wt% CNC loadings with cross-linking agents, formation of the ester-linkages led to significant decrease in permeability (by about 50%) as well as improvement in strength and toughness (by about 150%) of the processed films. The decrease in water vapor transmission rates is attributed to the reduction in the solubility of water molecules into PVA/CNC films due to the presence of the ester cross-links [87]. Similarly, decrease in both oxygen and WVP was observed when CNCs with distinct rod like morphology and aspect ratio of 20 to 35, was dispersed into biodegradable polymers PLA and PHB [19] through simple solvent exchange-cum-solution casting approach [42, 81]. Moreover, polymorphism of CNCs (cellulose I and cellulose II) leads to variation in its physico-chemical properties and morphology which, in turn, significantly alters the permeation behavior of molecules across the membranes. CNCs with cellulose II polymorphism have fibrous morphology with enhanced inter-molecular hydrogen bonding with PLA matrix, which significantly reduces both water vapor and oxygen permeation by about 50% and 80%, respectively [42]. In a recent study by these authors, it was found that surface of the CNCs could be modified with different anionic moieties even during the acid hydrolysis step. In addition, modification of CNCs with iron oxide nanoparticles resulted in magneto-stimuli responsive characteristics [88, 77]. The orientation of magnetic CNCs through directional alignment of magnetic field leads to significant orientation of CNCs in perpendicular or parallel directions through non-invasive magnetic field and may lead to improved barrier properties (as shown in Figure  12.4) which can be finely tuned depending upon the targeted applications [84]. Similarly reduced water vapor and oxygen permeability (OP) were observed when CNCs were used as a nano-filler in other biodegradable polymers such as PCL [89], hydroxypropyl methylcellulose (HPMC) [90], gelatin [91], etc. Recently CNCs have been directly used as coating material on different substrates such as paper or other plastics through LbL approach. The CNC coating on paper improves the intrinsic permeability by over 90% along with drastic improvement in burst index and bending stiffness values [92]. This shows that CNCs can either be used as nano-filler in the fabrication of biodegradable films or as coatings, which can improve the shelf-life of the stored food items.

338 Advanced Green Composites

12.2.3 Chitosan Based Composites: Chemical Modifications, Properties Evaluation and Applications Chitin and chitosan (CH) are promising, versatile and highly bioactive polysaccharides, which exist in the category of biomaterials. Chitin, the second most abundant polysaccharide after cellulose produced by nature, is a nitrogenous polysaccharide, inelastic, white in color and is extracted from the internal structure and exoskeleton of invertebrates such as shellfish, lobsters, insects and many others. Primarily, it is a composition of 2-amino-2-deoxy-D-glucose units, which is a derivative of cellulose. The only difference between cellulose and chitin is the presence of acetamide (-NHCOCH3) functional group in place of -OH group at C2 position. Approximately, ten billion tons of chitin is produced every year from crustaceans, molluscs, algae, insects and fungus. It is biocompatible, renewable, biofunctional, biodegradable environmentally friendly, and acts as a chelating agent, drug carrier, water treatment additive, wound-healing agents, etc. [93]. However, chitin has limited applications in pharmaceutical industries as it exhibits water insolubility and extreme hydrophobic character due to its highly expanded hydrogen-bonded semicrystalline structure. Chitosan is a derivative of chitin and is produced by a simple step of deacetylation or depolymerization under the influence of strong alkaline conditions (concentrated sodium hydroxide) or enzymatic hydrolysis with the help of chitin deacetylase. Degree of deacetylation (DD%) is defined as the percentage removal of acetyl groups from chitin molecule and their conversion to amino groups [94]. Chitosan, β-(1,4)-2-amino-2-deoxy-D-glucopyranose, is a biodegradable and biocompostable polymer and its structure strictly depends on the DD%. When the DD% is lower than 50%, the material is still termed as chitin otherwise it is called chitosan. The extracted chitin from various resources contains distinct polymorphism in α, β and γ forms. The α and β forms of polymorphism represent the antiparallel and parallel arrangements of polymer chains, respectively, and random arrangement of polymer chains is termed as γ form. Chitosan shows enormous possibilities of physical and chemical modifications via ionic interactions and grafting due to the presence of -OH and -NH2 groups, which form a range of chitosan derivatives. In this section, comprehensive research on chitin and chitosan have been discussed based on their structures, physicochemical properties and applications [95]. In recent years, chitin and chitosan have been produced at commercial scale in many countries including India, Poland, Japan, Norway, Australia and USA. Both, chitin and chitosan are employed in countless applications e.g. as adsorbents for the removal of dyes [96]

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and metal ions [97] in waste water treatment, fat binder in food industry [98], thin membranes in filtration processes [99], edible layer in packaging [100], coating material for fertilizer and seeds in agriculture [101], food additive and preservative [102], controlled release of agrochemicals [103], lotions and body creams in cosmetics [104], adhesive paper [105] and surface treatment [106] in pulp and paper industry, wound healing [107], tissue engineering [108], gene delivery [109], vehicle for drug delivery [110] as shown in Figure 12.5. Such significant contribution in various applications has been possible only because of its versatile properties. Chitosan is a nitrogen rich linear amino-polysaccharide with rigid structure of D-glucosamine which shows high hydrophilicity and low crystallinity as compared to that of chitin. The amino group of chitosan gets deprotonated at low pH and behaves like a powerful nucleophile which makes it soluble in acidic environment and facilitates gel like structure due to intermolecular hydrogen bonding. The chemical activation and cross-linking of chitosan with other polymers are easily possible due to the presence of reactive groups. Chitosan is a polyelectrolyte, non-toxic, antimicrobial, antifungal, antibacterial, antiulcer and blood anticoagulant [111]. The extraction of chitin is possible only from the cuticles of marine creatures and insects such as crustaceans, lobster, shrimps, crab, cockroaches, worms, mushrooms, etc. Chitin is available as a constituent of shell composition along with the complex network structure among proteins,

Food preservation from microbial deterioration Waste material recovery from food processing discards

Emulsifying agent Animal feed additive

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Enzyme and cell immobilization Cell stimulating materials

Water purification, sludge treatment Reverse osmosis, filtration membranes Production of biodegradable packaging films

Figure 12.5 Potential applications of chitosan and its derivatives in various sectors.

340 Advanced Green Composites polysaccharides and calcium carbonate. Hence, the extraction of chitin from renewable resources involves two major techniques; chemical extraction and biological extraction. In the chemical extraction process, the removal of proteins and calcium carbonate is performed by deproteinization and demineralization, respectively. A small amount of lipids and pigments that are also present in the cuticles are removed with the help of decoloring agents such as acetone, hydrogen peroxide, etc. [112]. Biological extraction is advantageous over chemical extraction, as it prevents the deterioration in physico-chemical properties. Chitosan is widely used in packaging applications due to its biodegradability, non-toxicity and film forming ability. This property was utilized by Khan et al., who fabricated biodegradable films using chitosan as matrix and nano-crystalline cellulose (NCC) as filler with 1-10 wt% loading, using the solution casting method [113]. An improvement of about 26% was noticed in ultimate tensile strength (UTS) as compared to that of pristine chitosan film with 5 wt% loading of NCC. Chitosan/NCC films demonstrated an optimum reduction of 27% in WVP as compared to that of neat chitosan, which is a requirement for packaging [113]. Chitosan was also used to prepare chitosan/tripolyphosphate nanoparticles (CH-TPP) by ionic gelation method and blended with HPMC to prepare edible films by solution casting method [114]. The available pores in HPMC matrix were collapsed by adding chitosan nanoparticles, which improved WVP, tensile properties and thermal properties. The WVP of HPMC films was reduced by upto 26% after the incorporation of CH-TPP nanoparticles [114]. Other examples for property enhancement by adding chitosan are grease-proof paper with chitosan coating [115], pulp fiber-chitosan sheets [116], biodegradable rice starch/chitosan films [117], chitosan/layered silicate nanocomposite films [118], chitosan-starch films loaded with silver nanoparticles [119] and corn starch/chitosan composite films [120]. In all cases mentioned above, it was observed that the presence of chitosan enhanced the barrier and other properties of various conventional and biobased polymers. While the physico-chemical properties of polymers are improved by the addition of chitosan, the hydrophilic nature of chitosan is a real drawback, which inhibits its contribution in packaging. Hence, the modification of existing biobased polymers has attracted much attention of researchers to overcome its limitations and it can be achieved by using many different techniques including copolymerization, grafting, grafting-cum-condensation polymerization or ring opening polymerization, etc. Among various modification techniques used, grafting, which can be achieved using various intermediates including free radical, ions, photo-initiated grafting, plasma induced grafting and enzymatic grafting, has received significant attention [121]. Photochemical

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grafting can be achieved by two ways; with sensitizer or without sensitizer [122]. The newest developed technique is enzymatic grafting, which performs with the help of enzymes. In this method, the enzymes behave as the initiator to perform the chemical grafting. For example, tyrosinase has the capability to transform phenol into o-quinone, which further reacts with chitosan under non-enzymatic conditions and forms grafted copolymer [123]. Different routes of surface modification of chitosan are shown in Figure 12.6. There are many advantages of grafting, including the transformation in hydrophilicity/hydrophobicity of polymers, enhanced thermal stability, improved mechanical and barrier properties. Chitosan is extensively used to fabricate tough, clean and flexible films with low OP and can be used in the form of edible packaging to enhance the food shelf life. It is noteworthy to mention that chitosan can’t be utilized on its own for packaging due to its hydrophilic nature and its incompatibility with hydrophobic polymers such as PLA leads to phase separation which reduces its mechanical and thermal properties [124]. Chitosan can be easily modified through the hydroxyl and amino groups to overcome the limitation of compatibility with PLA and other hydrophobic polymers. Several

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Figure 12.6 Different routes for surface modification of chitosan.

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342 Advanced Green Composites attempts have been made to modify chitosan with other conventional and biobased polymers, by using catalysts/initiators, which work as a bridge between two polymers [125-127]. Li et al., [128] and Zhou et al., [129] developed an amphiphilic graft copolymer using hydrophilic chitosan and hydrophobic poly (L-lactic acid) (PLLA) by protection-graft-deprotection mechanism. In this reaction mechanism, the monomer unit of PLLA reacts with the available hydroxyl groups on phthaloylchitosan (PHCS) in the presence of homogeneous system of N,N-dimethylformamide (DMF) solvent. The grafting percentage depends on reaction time and the PLLA/ PHCS ratio, which was varied between 43.4 and 342.6%. The grafted copolymer disperses in the form of solid and hollow spherical micelles in aqueous solutions as confirmed by TEM analysis and should be nontoxic due the use of biodegradable reactants. Chitosan has also been modified with PLA to form PLA grafted chitosan (PLA-g-CH) nanoparticles to load the drug, amphotericin B, by dialysis method for ocular delivery [129]. It was projected that the therapeutic gaps in ocular drug delivery can be conquered with the help of distinct structure of PLA-g-CH [129]. Similarly, Li et al. developed a chemical conjugation method for the synthesis of chitosan grafted PLA (CH-g-PLA) with low degree of grafting by using 1-ethyl-3-(3-dimethylaminopropyl) carbodiimide (EDC) and N-hydroxy-succinimide (NHS) as mediators for coupling reaction as shown in Figure  12.6 [126]. The molecular weight and amount of PLA can easily tune the chain length of grafted polymer. The grafted copolymer (CH-g-PLA) exhibits tunable hydrophilicity, hydrophobicity, solubility, thermal stability, biodegradability and self-assembling behavior, which can be utilized in various applications such as food packaging, drug delivery and tissue engineering. The biodegradable chitosan is also grafted with other polymers such as poly(p-dioxanone) (PDO) to form chitosan-graftpoly(p-dioxanone) copolymer (CGP) in the presence of a catalyst, stannous octoate (SnOct2), using ring opening polymerization, which is highly used in drug delivery [127]. Similarly, Schreiber et al., studied the antioxidant activity of pristine chitosan and gallic acid grafted chitosan films for their use in food packaging application [130]. The grafting of gallic acid on chitosan backbone was performed in the presence of EDC and NHS. An improvement in 2,2-diphenyl-1-picrylhydrazyl (DPPH) scavenging of 89.5% was observed in gallic acid grafted chitosan films in comparison to that of pristine chitosan, which is highly appreciable for its use in multifunctional packaging applications [130]. Chitosan can be modified not only by high molecular weight polymers but also with low molecular weight polymers such as oligo (L-lactic acid) (OLLA) based on its use in particular application. Yao et al., [131] synthesized low molecular weight

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chitosan grafted oligo (L-lactic acid) copolymers (CH-g-OLLA) with the help of 1,6-hexanediol, dimethylsulfoxide (DMSO) and acetic anhydride. The molecular weight of CH-g-OLLA varied from about 600 to 5000 Da with the variations in side chain length. The CH-g-OLLA copolymer can be widely used in packaging, tissue engineering and drug delivery [131]. Further, Wu et al., also prepared polymeric micelles of graft copolymer based on amphiphilic chitosan and DL-lactide in the presence of triethylamine and DMSO (as shown in Figure 12.6), which can be an appropriate option for improving gas barrier properties of packaging materials, entrapment of drugs and their controlled release [132]. From the above discussion, it is quite clear that biodegradable chitosan can be easily modified with various polymers in the presence of different initiators by using a number of techniques. Although, these polymers are biodegradable and non-toxic, the initiators are hazardous and toxic in nature and are not acceptable for packaging and medical applications. Hence, Pal and Katiyar synthesized low molecular weight lactic acid oligomer-grafted-chitosan (OLLA-g-CH) bionanocomposites by condensation polymerization reaction in microwave without using any catalyst/initiator as shown in Figure  12.6 [25]. OLLA-g-CH is amphiphilic in nature and behaves like self-assembled polymeric micelles that are uniformly dispersed in PLA matrix and form PLA-chitosan bionanocomposite films. Such films are highly recommended for food packaging applications due to their improved OP by up to 10 folds as compared to pure PLA. Fabrication of PLA/OLLA-g-CH bionanocomposite films is shown in Figure  12.7. This significant reduction is observed due to the reduction of oxygen solubility in the presence of filler and elevation in crystal nucleation density [25]. Similarly, various researchers have utilized chitosan with other distinct polymers to enhance the physico-chemical and barrier properties. From the above discussion it can be concluded that chitosan is a versatile, biodegradable, non-toxic, biocompostable biopolymer and is highly recommended for various applications.

12.2.4 Natural Gum Based Composites: Chemical Modification, Property Evaluation and Applications The natural gums are non-toxic, biodegradable and low cost polysaccharides, which offer many benefits over fossil based synthetic polymers due to their stabilizing and emulsifying properties. As a result, they are applicable in various areas such as pharmaceuticals, adhesives, films, paints, beverages, and cosmetics [133]. Gum arabic (GA) is one of the natural complex polysaccharides, derived from an exudate of Acacia trees namely Acacia

344 Advanced Green Composites

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Figure 12.7 Fabrication of PLA/OLLA-g-CH bionanocomposite films.

Senegal [134]. It is structurally composed of three main components i.e. arabinogalactan (AG, ~90%), arabinogalactan-protein (AGP, ~10%) and glycoprotein (GP, ~1%), which are mainly different in molecular weight and protein content [135]. Structural fragments of GA and the various components are shown in Figure  12.8 (a) [135]. The backbone of GA is galactose which is highly branched with β-1,3-glycoside. The fractions of side chains vary from product to product and are composed of L-arabinose (24-29%), L-rhamnose (12-14%), D-galactose and D-glucuronic acid (1617%) [136]. The other existing gums are guar gum (GG), xanthan gum (XG) and gum rosin (GR). Guar gum is chemically termed as polygalactomannan, collected from seeds of leguminacea plant. Primarily, it is a heteropolysaccharide and is composed of mannose, galactose and sugars. The backbone of GG is a linear chain of (1-4)-linked β-D-mannopyranose with (1-6)-linked

Green Composites with Excellent Barrier Properties R

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Figure 12.8 (a) Structural fragments of gum arabic having various components i.e. arabinopyranose (Ap) [135]; rhamnose (R); arabinose (A); galactose (G); 4-O-methyglucuronic acid (Um) and glucuronic acid (U), chemical structural of (b) guar gum [137]and (c) xanthan gum [138].

α-D-galactopyranose short side chains connected at every alternate mannose units as shown in Figure  12.8(b) [137]. It is hydrophilic in nature and able to absorb moisture due to the presence of hydroxyl groups. As a result, it can’t be used in paper, cosmetic and food industry. Xanthan gum is an exopolysaccharide produced by the Bacterium Xanthomonas campestris. It is an anionic hetero-polysaccharide with a primary structure consisting of repeated pentasaccharide units. The pentasaccharides are formed by two glucose units, two mannose units and one glucuronic acid unit in the molar ratio of 2.8:2.0:2.0 as shown in Figure 12.8(c) [138]. The

346 Advanced Green Composites primary structure of xanthan gum contains a cellulosic backbone (β-Dglucose residues) and a trisaccharide side chain of β-D-mannose-β-Dglucuronic acid-α-D-mannose attached to alternate glucose residues of the main chain. The non-terminal D-mannose unit in the side chain contains an acetyl functional group. The anionic character of this polysaccharide is due to the presence of both glucuronic acid and pyruvic acid groups in the side chain [139]. The researchers have used all above explained gums for various applications such as adhesive, packaging, food additive and medical fields. However, it is very difficult to disperse hydrophilic gums into hydrophobic polymers. The modification of gums using various grafting techniques can provide an alternate solution for the existing limitation. A copolymer of GA and polyaniline has been synthesized using peroxydisulfate as initiator and oxidant by radical polymerization as mentioned in Figure  12.9(a) [140]. Further, the above technique has been modified by Tiwari and Singh, 2008 and synthesized gum acacia-graftpolyaniline in the presence of ammonium peroxydisulfate as initiator and hydrochloric acid via microwave accelerated oxidative radical polymerization as shown in Figure 12.9(a) [141]. In the same line, Kaith et al., 2010 [142] have developed three-dimensional cross-linked structure with the help of grafting of GA with methacrylic acid in the presence of potassium persulfate as initiator and hexamethylene tetramine as cross-linker using free radical polymerization. It was observed that the copolymer produced by this method was sensitive to temperature and pH and also exhibited salt-resistant swelling. Pandey and Mishra used xanthan gum and ethylacrylate in the presence of potassium peroxydisulphate (KPS) as initiator to synthesize xanthangraft-polyethylacrylate copolymer by microwave induced emulsion copolymerization [143]. The xanthan-graft-polyethylacrylate copolymer had the adsorption capacity of Pb2+ four times than that of pure xanthan gum. However, natural gums are less used in packaging applications due to their hydrophilic and adhesive behavior. Wyasu and Okereke 2012 [144] have been able to improve the film forming ability of GA by using some additives such as pure glycerol, diethyl glycol and ethylene glycol for their application in packaging. Peanut protein isolate (PPI) films were fabricated which were glycated with GA and compared with neat PPI and PPI-GA mixture films. The cross-linking of PPI with GA resulted in 2-2.5 fold reduction in WVP compared to neat PPI films. However, after 3 days of glycation, increase in tensile strength was observed with decrease in WVP. Prepared chitosan-GA (CH-GA) polyelectrolyte complex films were prepared by Sakloetsakun et al. to study the physicochemical, mechanical and mucoadhesive properties [145]. High mechanical

Green Composites with Excellent Barrier Properties

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Figure 12.9 Synthesis of (a) gum arabic-graft-polyaniline copolymer by radical polymerization and microwave-accelerated oxidative radical polymerization and (b) LA-g-GA by polycondensation reaction [56].

strength was observed for CH-GA films as compared to unmodified CH films. However, it was also experienced that the addition of GA reduces the extensibility of the films because of the cross-linking. Furthermore, no significant change in WVP was observed for CH-GA films. It was suggested that they interact via intermolecular hydrogen bonding and electrostatic forces due to which relatively lower WVP, low water uptake, erosion and puncture strength were observed as compared to neat chitosan films [145]. GA can also be used as edible coating due to its biocompatibility and nontoxic nature. The development of edible coating of GA with glycerol monostearate for tomato fruit enhances the shelf life and postharvest quality. The results indicated that GA delayed the ripening process and extending the storage life of tomatoes by up to 20 days without any off-flavor or spoilage [146]. Natural gums are not only used as edible coatings but also can be used as a filler with various polymers in food packaging. Onyari et al. prepared hydrophobic PLA and hydrophilic GA blend films using chloroform solution by conventional solvent casting method and conducted biodegradation study using thermophilic bacteria with higher efficacy [147]. Based on the literature survey, it can be found that natural gums have been modified using conventional additives that are toxic in nature and, in turn, create unwanted environmental problems. In order to overcome such problems, Tripathi and Katiyar

348 Advanced Green Composites synthesized lactic acid-grafted-gum arabic (LA-g-GA) using lactic acid and GA (as shown in Figure 12.9 (b)) without any external additive by polycondensation reaction under microwave radiation and subsequently LA-g-GA are used as a filler in PLA matrix to fabricate compatible PLA/ LA-g-GA bionanocomposite films by both solution casting and melt extrusion technique [56]. The glass transition temperature of bionanocomposite films was significantly reduced due to the plasticization effect. This, in fact, is a positive change in order to use this material for flexible packaging. The water vapor transmission rate and oxygen transmission rate (OTR) of PLA/LA-g-GA bionanocomposite films were reduced by ~27 and 90%, respectively, with an increase in filler concentration. Such a large reduction in OTR was only possible by the significant reduction in the solubility of oxygen in bionanocomposite films. Therefore, from the above discussion it can be inferred that utilization of the modified gums as fillers, through strategic chemical modifications by improving its dispersion in the polymeric matrix which leads to the development of biodegradable films with improved barrier properties.

12.2.5 Silk Based Composites: Property Evaluation, Chemical Modifications and Applications The term silk refers to the continuous filaments spun by species belonging to the family of Lepidoptera and Araeneae, the taxonomic classification can be as seen in Figure 12.10 a [148]. The silk fluid is produced by specialized glands and stored in the lumen sac before it is spun into fibers. The fibers have a range of purposes from cocoons which protect eggs/pupa to dragline silk by which the spiders hang and webs which can capture prey and some that are used for escape [149]. Domesticated silkworm variety such as Bombyx Mori spins dual-brin (filament) silk fibers made of fibroin (protein) on the inside and an outside protective coating of sericin (protein) with gum-like properties which bind the silk fibers in shape of cocoon [150] as seen in the sketch shown in Figure 12.10 b. Other than domesticated silkworm varieties there are other wild silk varieties such as tassar (tussah), eri and muga silk as seen in Figure 12.10c-e. Fibroin imparts the strength for which silk is known for whereas sericin is a hydrophilic amorphous gum-like coating which amounts to 25% of weight of cocoon [151]. Sericin is known to have antibacterial, UV cutting and moisture regulating properties making it a suitable choice for present day cosmetic ingredient [152]. Spider silk or spidroin cannot be produced on large scale because of the cannibalistic nature of spiders. Alternatively, genetic engineering has been used to produce recombinant spider silk.

Green Composites with Excellent Barrier Properties

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Figure 12.10 (a) Taxonomix classification of silk spinning organisms showing common varieties of cocoons available; (b) Mulberry silk (bombyx mori); structural arrangement of silk fiber of bombyx mori; (c) Eri silk (samia ricini); (d) Muga silk (antherea assama) [150].

Both silkworms and spiders employ silk spinning processes that are quite similar. However, it is still a mystery to researchers so as to what exactly happens on the molecular scale in the spinning duct, which makes the silk present in the sac in liquid form convert into fibers with high strength. Both types of silks do not have any common lineage and their amino acid composition is different as well. Fibrous proteins such as the silks and keratin generally have a highly repetitive primary sequence for example mulberry silkworm fibroin contains poly-Gly-Ala (glycine-alanine) while spiders and other wild silkworm fibroin contains blocks of poly-Ala forms β-sheet secondary structures called β-sheet crystallites. These are inextensible structures that are embedded in an amorphous extensible matrix in the form of random coils as shown in Figure  12.11  h [153]. Crystalline region in silkworm silk (40-50%) is higher as compared to that in spider silk (35%) [31]. Among all available biopolymers, silk is known for its strength, slow degradation, biocompatibility and non-toxicity. Mechanical properties of silk along with various other synthetic materials are presented in Table 12.1. Highly crystalline structure of silk fiber provides tensile strength to the

350 Advanced Green Composites 15–20 m

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Figure 12.11 Hierachical arrangement of silk (a) raw silk fiber, (b) dimesnions of silk fiber, (c) incompletely degummed silk fiber, (d) completely degummed silk fiber, (e) opical micrograph of degummed silk, (f) microfibrils emerging out of silk fibers upon size reduction, (g) nanofibrils emerging out of silk fibers upon size reduction [153, 157].

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silk, while high amount of hydrogen bonding makes it insoluble in most solvents. Silk has an orthorhombic unit cell structure and very high packing density which makes it difficult for any solvent such as dilute acids, dilute alkalis and water to enter the structure and disrupt it [154]. Various techniques including XRD, TEM, and NMR have been used to explore the silk structure [155]. The x-ray diffraction spectra confirm the presence of highly crystalline β-sheet structures made from highly repeating primary sequences like poly  Gly-Ala or poly-Ala. Where toxicity and incompatibility to biological systems limits the use of contemporary materials, silk can offer the much needed shift towards sustainable and active/smart packaging. As per the report published in 2009, silk with a world annual production of 140,000 tons amounts to just 0.4% share of the total estimated production of natural fibers [156]. Silk is mainly used for textiles while a very small share is diverted for other applications. Before using in any application, silk fibroin (silk fiber) needs to be isolated by removal of sericin layer and salt crystals by boiling silk fibers in washing soda which is known as the degumming process. As seen in Figure 12.11a, the raw silk fiber consists of silk fibroin core and outer sericin layer, which needs to be removed and Figure 12.11c shows silk fibroin with incomletely removed sericin layer which can be seen as patches. Upon rinsing and drying degummed silk fibroin is obtained as seen in Figure 12.11d and 12.11e. For food packaging applications, silk fibroin can be processed using two different pathways. The first one is by mechanical attrition or chemical treatment, breaking silk into microfibrils as can be seen in Figure 12.11f. These microfibrils can be further broken into nanofibrils which are in the form of bundles which can be seen in Figure 12.11 g. Nanofibril bundle mean width ranges from 90–170 nm, while each nanofibril is ~5 nm in diameter [157]. The second pathway is resolubilization using chaotropic solvents for aqueous and solvent formulations. Once the silk is resolubilized then it is used to create different material forms (films, microspheres, electrospun nanofibers, foams/sponges and hydrogels, etc.) with morphology consisting of random coil molecular conformation. This needs to be converted to β-sheet structure upon annealing by water vapor/alcohol or by mechanical stretching. By changing the silk structure by annealing with alcohol, water vapor or exposing to temperature or radiation, increase the crystallinity can be obtained. The process can be modified to fine-tune barrier and mechanical properties of silk [158]. Mechanical properties of biopolymers can be changed substantially by using silk fibers as a reinforcing material. Core-sheath type of composite fibers can be obtained using spidroin/PLA system, where what stays

352 Advanced Green Composites at the core is decided by the concentration of spidroin. In a study, Hardy and Scheibel found that when spidroin concentration was >15%, spidroin formed the core and improved the tensile properties of the fibers [150]. In another study composite PLA/fibroin fibers were found to have improved breaking strength [159]. Cheung et al. used silk fibers as reinforcement in PLA, which showed an increase in elastic modulus and ductility by 40% and 53%, respectively [160]. One of the main applications of silk, due to its biocompatibility, is in wound healing where silk/PEO electrospun mats exhibit good morphological and functional properties with very fast degradation of 86% in 14 days with OTR and WVTR of 15,460 cm3/m2/ day and 2,160 gm/m2/day, respectively. Similar type of mats can be used to prepare breathable packaging which can protect fruits from flies, dust, direct air contact thus preventing discoloration, etc. [161]. Shanmugam and Sundaramoorthy prepared silk/PLA mats which showed a WVTR of 823  g/m2/day, it was interesting to note that neat silk film gave a higher WVTR of 1,323 g/m2/day [162]. These mats were targeted for wound healing applications. Perfectly crystalline composite films of fibroin/nylon 66 were solvent cast from formic acid solution by Liu et al. [163]. Other film forming techniques include spin-coating, compression molding or solvent based layer-by-layer assembly processes using silk. It is important to note that dry-cast silk films are brittle with low breaking strain ~0.5 to 3% limiting its use in flexible packaging [164]. On the other hand ultrathin films formed either by spin coating or layer-by-layer techniques have toughness of ~328 KJ.m-3 and ultimate tensile strength of about 100 MPa along with flexibility. These properties are similar to many currently available polymers [165]. Bio-nanocomposite films made of the amphiphilic and cationic recombinant spider silk protein as matrix and the synthetic layered silicate sodium hectorite (Na-hec) as filler were produced in an all-aqueous suspension casting process. In comparison to the uncoated PET foil (OTR: 22.83 ± 0.33 cm3 m−2 day−1 bar−1; WVTR: 4.20 ± 0.64 g m−2day−1), the OTR of bionanocomposite-coated PET was reduced by 90% to 2.32 ± 0.38  cm3  m−2 day−1 bar−1, while the WVTR decreased by 96% to 0.18 ± 0.05 g m−2day−1 [166]. The large reduction of WVTR was remarkable, and barrier coatings with the detected properties could find application in the field of food packaging [166]. Biopolymer nanocomposite films were prepared by adding GO and reduced GO (RGO) to sodium carboxy-methyl cellulose (CMC)/silk fibroin (SF) matrix. Thermal stability, glass transition temperature and surface roughness of the nanocomposite films were gradually increased with the incorporation of RGO. In particular, an increased beta-sheet content corresponds to a reduction in

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oxygen diffusion through silk fibroin thin films [103]. Simple dip coating of fruits in silk fibroin resulted in formation of micrometer thin silk films, which needed to be annealed. This conformation change decreased the by 50 folds to 2 10-11 cm2s and water vapor permeability to 6.49 10−11 g m−1s−1Pa−1 [167]. These only silk containing films can replace the cling film and might be very useful later. Composite films of PVA/RS/ SF showed high strength, high flexibility, transparency, low water swelling and transparency when PVA/RS were mixed in the ratio of 60:40 with 2% silk addition. It was observed that high content of silk increased the OP[168]. Fibroin and PLA both are soluble in dioxane and interact through hydrogen bonding. When PLA was added to more than 10 wt.% the composite was found to have increased tensile strength and elongation [169]. Poly (allylamine) is readily soluble in water but when composite films of fibroin and poly (allylamine) were prepared, they were found to be more stable in water. This study suggests that addition of fibroin to the polymer hampers the wetting capacity which is highly required for food packaging [170]. Whey protein isolates blended with silk sericin also rendered films that had increased mechanical properties and low water vapor permeation and, in addition, they were both edible and biodegradable [171]. Recently, the naturally phase-separated arrangement was artificially recreated in a material composed of a laminate of chitosan and fibroin layers. This composite has been called “Shrilk” and is twice as strong as its strongest component (chitosan) and ten times stronger than the simple blend of both components. The fact that chitosan and fibroin are already approved by the United States Food and Drug Administration (FDA) for medical use in humans also opens up the possibility of their use as food packaging films [172]. Similarly, silk and carrageenan have been used to prepare antimicrobial films to store sausages [173]. Gu et al. prepared SF/chitosan films by solvent casting and obtained a water vapor transmission rate of 2500 g/m2/day [174]. In their studies, addition of silk was seen to lower the WVTR [174]. In another study silk sericin was added to the collagen matrix and water vapor transmission rate was improved to 1013 from 980 gm/m2/day, which was targeted for wound healing applications [175]. A number of living organisms have the ability to biosynthesize polymer crystals [176]. In the case of silk, the water insoluble and highly crystalline residue obtained after removing the amorphous bulky and polar amino acids is referred to as silk crystals. Silk nanocrystals could be quite promising when compared to inorganic nanocrystals. They are obtained from nature and have the following advantages such as being stable in suspension, easily available, inexpensive, renewable, easy chemical modification,

354 Advanced Green Composites mechanical stability to withstand processing conditions, long-term sustained release of drugs and slow degradability. Nova et al., [177] used molecular modelling approach to study the interplay between crystalline beta sheet motifs and non-crystalline regions of silk. They observed that the amorphous regions unravel first imparting the silk its extensibility while the beta sheet nanocrystals provide the ultimate tensile strength of silk, depending upon the size of crystals [177]. Silk fibroin nanocrystals have the originality of the platelet-like morphology [178], and consist of crystalline nanoplatelets having a diameter