Fire Retardancy of Polymers: New Strategies and Mechanisms

Fire Retardancy of Polymers New Strategies and Mechanisms

This book is dedicated to our families, for the support and tolerance they showed during its production, particularly to Helen, Sam, Matthew, Isobel, Baldev, Abhineet, Navrohit and Amitarun.

Fire Retardancy of Polymers New Strategies and Mechanisms

Edited by T Richard Hull Centre for Fire and Hazard Sciences, University of Central Lancashire, Preston, UK

Baljinder K Kandola Centre for Materials Research and Innovation, The University of Bolton, Bolton, UK

ISBN: 978-0-85404-149-7 A catalogue record for this book is available from the British Library r Royal Society of Chemistry 2009 All rights reserved Apart from fair dealing for the purposes of research for non-commercial purposes or for private study, criticism or review, as permitted under the Copyright, Designs and Patents Act 1988 and the Copyright and Related Rights Regulations 2003, this publication may not be reproduced, stored or transmitted, in any form or by any means, without the prior permission in writing of The Royal Society of Chemistry or the copyright owner, or in the case of reproduction in accordance with the terms of licences issued by the Copyright Licensing Agency in the UK, or in accordance with the terms of the licences issued by the appropriate Reproduction Rights Organization outside the UK. Enquiries concerning reproduction outside the terms stated here should be sent to The Royal Society of Chemistry at the address printed on this page. Published by The Royal Society of Chemistry, Thomas Graham House, Science Park, Milton Road, Cambridge CB4 0WF, UK Registered Charity Number 207890 For further information see our web site at www.rsc.org

Preface This volume follows in a tradition of ten previous biennial meetings and publications by those working to reduce the hazards associated with burning polymers. The 11th Meeting on Fire Retardant Polymers (FRPM’07), which took place in the Albert Halls in Bolton in July 2007, brought together over 200 scientists from across the globe with representatives from over 20 countries, to discuss the latest developments in fire retardant technology. There were around 100 presentations at the meeting, and as editors we have extracted the highlights, summarised in 25 chapters, in order to present the state of the art in fire retardant strategies and mechanisms. All of the papers in this volume were subject to peer review, by two independent experts in fire retardancy. We are extremely grateful for the efforts put in, both by our authors and by the reviewers in helping us to maintain the highest standards in this volume. The volume opens with a brief overview on Polymers and Fire setting the context by outlining the stages of a fire and the broad groups of fire retardants and their modes of action. The first section discusses the new Fire Retardant Strategies which are currently emerging. The Synergy between nanometric alumina and organoclay in conventional fire retardant systems for ethylene-vinyl acetate leads in with the novel approach of investigating the effect of conventional fire retardants in nanoscopic form. Strained organophosphorus compounds as reactive flame retardants for polymeric materials continues the quest for halogen free fire retardants by designing phosphorus compounds which will decompose in sequence with the polymer to optimise the flame retardant effect. Amorphous silicon dioxide as additive to improve the fire retardancy of polyamides describes the novel application of smooth silica microspheres for flammability reduction. Use of organosilicone composites as flame retardant additives and coatings for polypropylene demonstrates the dual effects on flame retardancy and reducing surface adhesion of treatments based on polyborosiloxane. Organo-modified Fire Retardancy of Polymers: New Strategies and Mechanisms Edited by T Richard Hull and Baljinder K Kandola r Royal Society of Chemistry 2009 Published by the Royal Society of Chemistry, www.rsc.org

v

vi

Preface

ultrafine kaolin for mechanical reinforcement and improved flame retardancy of recycled PET shows how the interaction of ultrafine kaolins with triphenylphosphite leads to the formation of a barrier which enhances charring. Complex micro-analysis assisted design of fire retardant nanocomposites-contribution to the nano-mechanism uses a novel laser-Raman system to simultaneously decompose and probe the decomposition products of fire retardant polymer systems. The second section is focussed on Nanoparticulate Fillers which have combined the potential to improve the physical properties of polymeric materials while also reducing their flammability. The section opens with a discussion on the Impact of nano-particle shape on the flammability of nanocomposites. The majority of fire retardant nanotechnology is currently focussed on the incorporation of nanoscale carbon and clay to form polymer nanocomposites. The Thermal and combustion behaviour of polymer-carbon nanofibre composites describes an approach with one form of these new materials. This is followed by an extensive investigation of the Combination of carbon nanotubes with fire retardants: the thermal and fire properties of polystyrene nanocomposites. Significant assessment of nanocomposites combustion behaviour by the proper use of the cone calorimeter describes the interpretation of the novel burning behaviour of polymer nanocomposites, and its investigation using cone calorimetry. The combination of traditional and nanocomposite approaches to fire retardancy using Phosphorus based epoxy resins nano-clay composites is reported. A novel approach to the somewhat elusive, and variable, mechanisms of nanocomposite fire retardancy is approached using viscosity measurements in the Study of the relationship between flammability and melt rheological properties of flame retarded poly(butylene terephthalate) containing nanoclays. The section concludes with another investigation combining novel and nanocomposite fire retardants outlining the Thermal and fire performance of flame-retarded epoxy resin: Investigating interaction between resorcinol bis(diphenyl phosphate) and epoxy nanocomposites. The third section investigates another important class of fire retardants, Intumescents, whose effect is created by the formation of a barrier layer which is then driven from the surface of the substrate by the release of a volatile or gaseous component. These may take the form of coatings or may be incorporated into the bulk of the material. A major application is in the creation of a polymer-based coating which is applied to structural members such as steelwork for fire protection. The section opens with an analysis of the bubble formation responsible for blocking the transfer of heat using Image analysis of 2D intumescent char sections to estimate porosity. This is followed by a description of Efficient modelling of temperatures in steel plates protected by intumescent coating in fire. The section finishes with a discussion focussed around Fire retardancy and fire protection of materials using intumescent coatings - a versatile solution? The fourth section describes one of the major and most challenging application areas for fre retardant systems, namely Fibres and Textiles. The section opens with a discussion of Trends in textile flame retardants – a market review.

Preface

vii

This is followed by an overview of New and potential textile flammability regulations and test methods within the USA. From this contextual setting of the market and regulatory framework, an investigation of the Flame retardancy of cellulosic fabrics: interactions between nitrogen additives and phosphoruscontaining flame retardants is described. The section then covers two detailed investigations into the fire retardancy of polyacrylonitrile fibres. The Synergistic flame retardant copolymeric polyacryonitrile fibres containing dispersed phyllosilicate clays and ammonium polyphosphate is reported, followed by reactive fire retardant approach to Flame retardance of polyacrylonitriles covalently modified with phosphorus- and nitrogen-containing groups. This is followed by Novel fire retardant back-coatings for textiles. The section concludes with a description of The effect of yarn and fabric construction and colour in respect of red reflectance and pigmentation on the thermal properties and LOI of flame retardant polypropylene fabrics. The final section of the book describes investigations into the effects of fire retardant on the major hazard to life in fire, that resulting from Fire Toxicity. The Influence of fire retardants on toxic and environmental hazards from fires describes a large body of work assessing the hazards associated with burning polymers and their flame retardant counterparts. The volume concludes with a discussion on the specific influence of nanocomposite formation and fire retardants on the toxic product yields of nylon 6 and polypropylene in Assessment of fire toxicity from polymer nanocomposites. We trust that the book will provide fire retardant materials’ developers with latest developments in research or design of new fire retardant materials and understanding their mechanisms of action. Finally, we would like to thank our colleagues, Prof Dennis Price, Prof Richard Horrocks, Dr Shonali Nazare, Dr Everson Kandare and Dr Anna Stec for their support during completion of this book.

Contents Introduction

Polymers and Fire T.R. Hull and A.A. Stec 1 2

Hazards from Fire Fires and Fire Growth 2.1 Conditions of Each Fire Stage 2.2 Chemical and Physical Processes 2.3 Studying Polymer Decomposition 3 Fire Effluent Toxicity 4 Structural Deformation 5 Fire Retardant Strategies 5.1 Physical Action 5.2 Chemical Action 5.2 Polymer Nanocomposites 6 Conclusions References

1 2 5 6 7 8 9 9 10 10 11 12 14

Fire Retardant Strategies Chapter 1

Synergy between Nanometric Alumina and Organoclay in Conventional Fire Retardant Systems for Ethylene–Vinyl Acetate N. Cinausero, J.-M. Lopez-Cuesta, F. Laoutid, A. Piechaczyk and E. Leroy 1.1 1.2

Introduction Experimental 1.2.1 Materials 1.2.2 Processing 1.2.3 Testing

17 18 18 18 18

Fire Retardancy of Polymers: New Strategies and Mechanisms Edited by T Richard Hull and Baljinder K Kandola r Royal Society of Chemistry 2009 Published by the Royal Society of Chemistry, www.rsc.org

ix

x

Contents

1.3 Results and Discussion 1.4 Conclusion References Chapter 2

Strained Organophosphorus Compounds as Reactive Flame Retardants for Polymeric Materials Bob A. Howell 2.1 2.2

Introduction Experimental 2.2.1 Materials 2.2.2 Initiator, 2,4,4,5,5-Pentaphenyl-1,3,2-Dioxaphospholane 2.2.3 Polymers 2.3 Results and Discussion 2.3.1 Thermal Properties of Styrene Polymers Containing Phosphorus Units 2.3.2 Evaluation of Flammability 2.4 Conclusions References

Chapter 3

28 29 29 29 29 30 33 33 34 34

Amorphous Silicon Dioxide as Additive to Improve the Fire Retardancy of Polyamides G. Schmaucks, B. Friede, H. Schreiner and J.O. Roszinski 3.1 3.2

Introduction Experimental 3.2.1 Materials 3.2.2 Sample Preparation 3.2.3 Test Methods 3.3 Results and Discussion 3.4 Conclusion Acknowledgement References

Chapter 4

19 26 27

35 40 40 40 40 41 47 47 48

Use of Organosilicone Composites as Flame Retardant Additives and Coatings for Polypropylene + A. Szabo´, K. Kiss and G. Marosi B.B. Marosfoi, 4.1 Introduction 4.2. Experimental 4.2.1 Materials 4.2.2 Sample Preparation 4.2.3 Preparation of PP Compounds

49 50 50 51 51

xi

Contents

4.2.4

Preparation of Composites with Multilayer Structures 4.2.5 Characterisation 4.3 Results and Discussion 4.3.1 Thermo-Oxidative Stability 4.3.2 Combustion Characteristics of Polypropylene-Based Composites 4.3.3 Multilayer Structure PP–(pBSil–OSEP–MB) 4.4 Conclusion Acknowledgements References Chapter 5

53 56 57 57 58

Organomodified Ultrafine Kaolin for Mechanical Reinforcement and Improved Flame Retardancy of Recycled Polyethylene Terephthalate B. Swoboda, E. Leroy, J.-M. Lopez Cuesta, C. Artigo, C. Petter and C.H. Sampaio 5.1 5.2

Introduction Experimental 5.2.1 Materials 5.2.2 Processing 5.2.3 Characterization Techniques 5.3 Results and Discussion 5.3.1 Properties of Unmodified Kaolins 5.3.2 Grafting of TPP onto Kaolin Surface 5.3.3 Morphological, Rheological and Mechanical Properties of Polymer Compounds 5.3.4 Thermal Stability and Reaction to Fire of Polymeric Compounds 5.4 Conclusion 5.5 Acknowledgements References

Chapter 6

51 52 52 52

59 61 61 61 62 64 64 64

67 69 73 73 73

Complex Micro-analysis Assisted Design of Fire-Retardant Nanocomposites – Contribution to the Nanomechanism + P. Anna and Gy. Marosi A. Szabo´, B.B. Marosfoi, 6.1 6.2

Introduction Experimental 6.2.1 Materials 6.2.2 Methods

74 78 78 78

xii

Contents

6.3

Results and Discussion 6.3.1 Nanonetwork Formation 6.3.2 Intumescent Polymeric Particle Formation 6.4 Conclusion Acknowledgements References

80 80 84 88 89 89

Nanoparticulate Fillers Chapter 7

Impact of Nanoparticle Shape on the Flammability of Nanocomposites F. Yang, I. Bogdanova and G. L. Nelson 7.1 7.2

Introduction Experiment 7.2.1 Preparation of Polymer–Inorganic Nanocomposites 7.2.2 Mechanical Testing of Polymer–Inorganic Nanocomposites 7.2.3 Morphology Study for Polymer–Inorganic Nanocomposites 7.2.4 Thermal Degradation of Polymer–Inorganic Nanocomposites 7.2.5 Flammability of Polymer–Inorganic Nanocomposites 7.3 Results and Discussion 7.3.1 Polycarbonate–Inorganic Nanocomposites 7.3.2 PS–Inorganic Nanocomposites 7.4 Conclusion References

Chapter 8

95 96 96 96 96 97 97 97 97 102 107 108

Thermal and Combustion Behaviour of Polymer–Carbon Nanofibre Composites D. Tabuani, S. Pagliari, W. Gianelli and G. Camino 8.1 8.2

Introduction Materials and Methods 8.2.1 Melt Blending 8.2.2 From Solution 8.2.3 Characterization 8.3 Results and Discussion 8.3.1 Morphology 8.3.2 Thermal Behaviour 8.4 Conclusions References

110 111 111 111 112 113 113 115 123 123

xiii

Contents

Chapter 9

Combination of Carbon Nanotubes with Fire Retardants: Thermal and Fire Properties of Polystyrene Nanocomposites Florentina Tutunea and Charles A. Wilkie 9.1 9.2

Introduction Experimental 9.2.1 Materials 9.2.2 Preparation of Composites 9.2.3 Instrumentation 9.3 Results and Discussion 9.3.1 Thermogravimetric Analysis 9.3.2 Cone Calorimeter Evaluation 9.4 Conclusions References

Chapter 10

Significant Assessment of Nanocomposites’ Combustion Behaviour by the Appropriate Use of the Cone Calorimeter A. Fina, F. Canta A. Castrovinci and G. Camino 10.1 10.2

Introduction Experimental 10.2.1 Materials 10.2.2 Preparation and Characterization 10.2.3 Combustion Tests 10.3 Results and Discussion 10.3.1 8 mm Specimens Combustion Behaviour 10.3.2 16 mm Specimens Combustion Behaviour 10.4 Conclusions Acknowledgements References

Chapter 11

125 126 126 127 127 127 127 135 145 146

147 148 148 148 150 151 151 153 157 158 158

Phosphorus-Based Epoxy Resin–Nanoclay Composites Jianwei Hao, Yanbing Xiong and Na Wu 11.1 11.2

Introduction Experimental 11.2.1 Materials 11.2.2 Preparation of Phosphorus-Based Epoxide 11.2.3 Preparation of Phosphorus-Based Epoxy–Nanoclay Composites 11.2.4 Characterization 11.3 Results and Discussion

160 161 161 162 162 162 163

xiv

Contents

11.3.1

Structure of EPO–P Analysis with FTIR 11.3.2 Structure of EP–P–nano Composites Analysis with XRD and TEM 11.3.3 Combustion Performance and Mechanical Properties 11.4 Conclusion References Chapter 12

163 164 167 167

Study of the Relationship Between Flammability and Melt Rheological Properties of Flame-Retarded Poly(Butylene Terephthalate) Containing Nanoclays S. Nazare, T. R. Hull, B. Biswas, F. Samyn, S. Bourbigot, C. Jama, A. Castrovinci, A. Fina and G. Camino 12.1 12.2

Introduction Experimental 12.2.1 Materials 12.2.2 Sample Preparation 12.2.3 Characterization and Testing 12.3 Results and Discussion 12.3.1 Nanodispersion 12.3.2 Differential Scanning Calorimetry and Thermal Analysis 12.3.3 Melt Viscosity 12.3.4 Flammability 12.4 Conclusions Acknowledgements References Chapter 13

163

168 170 170 170 171 172 172 174 177 178 182 182 183

Thermal and Fire Performance of Flame-Retarded Epoxy Resin: Investigating Interaction Between Resorcinol Bis(Diphenyl Phosphate) and Epoxy Nanocomposites Charalampos Katsoulis, Everson Kandare and Baljinder K. Kandola 13.1 13.2

Introduction Experimental 13.2.1 Materials 13.2.2 Sample Preparation and Characterization

184 185 185 185

xv

Contents

13.2.3 13.2.4 13.3 Results 13.3.1 13.3.2

Thermogravimetric Analysis Flammability Tests and Discussion XRD and TEM Analysis Thermal Degradation Behaviour of Epoxy Resin and Its Composites 13.3.3 Flammability Behaviour 13.4 Conclusions Acknowledgements References

186 186 187 187 188 195 202 203 203

Intumescents Chapter 14

Porosity Estimates of Intumescent Chars by Image Analysis J.E.J. Staggs 14.1 14.2 14.3

Introduction Pore-Finding Algorithm Relationship Between Area Porosity and Volume Porosity 14.4 Relationship Between 2D and 3D Pore Distributions 14.4.1 Test Case 1 (Identical Spheres) 14.4.2 Test Case 2 (Spheres with Uniformly Distributed Radii) 14.5 Construction of 3D Distributions from 2D Distributions 14.6 Analysis of a Real Char Section 14.7 Conclusion Acknowledgements References Chapter 15

209 212 213 215 216 217 217 219 223 224 224

Efficient Modelling of Temperatures in Steel Plates Protected by Intumescent Coating in Fire J.F. Yuan and Y.C. Wang 15.1 Introduction 15.2 Mathematical Modelling 15.3 Validation 15.4 General Analysis of Intumescence Process 15.5 Parametric Studies 15.6 Conclusions Nomenclature References

225 227 230 231 233 238 238 239

xvi

Chapter 16

Contents

Fire Retardancy and Fire Protection of Materials using Intumescent Coatings – A Versatile Solution? S. Duquesne, M. Jimenez and S. Bourbigot 16.1 16.2

Introduction The Use of Intumescent Coatings for The Fire Protection of Steel Structures 16.3 Fire Protection of Polyurethane Foams using Intumescent Systems 16.4 Conclusion References

240 241 245 252 252

Fibres and Textiles Chapter 17

Trends in Textile Flame Retardants – a Market Review R. Hicklin, R. Padda and G. Lenotte 17.1 Introduction 17.2 Burning Behaviour of Cotton Fabrics 17.2.1 Factors that Affect the Burning Behaviour of Cotton Fabrics 17.2.2 Combustion of Cotton 17.3 Mechanism of Phosphorus Flame Retardants 17.4 Classification of Flame Retardants 17.5 Flame Retardant Selection 17.5.1 Non-Durable Flame Retardants 17.5.2 Semi-Durable Flame Retardants 17.5.3 Durable Flame Retardants 17.6 Flammability Standards and Testing 17.7 Health, Safety and Environmental Considerations 17.8 Fibre Blends 17.9 Future Developments 17.9.1 Multifunctional Fabrics 17.9.2 Alternative Chemistries 17.9.3 Application Technologies 17.10 Conclusions References

Chapter 18

255 256 256 256 256 257 257 257 259 259 261 261 262 263 263 263 263 264 264

New and Potential Textile Flammability Regulations and Test Methods within the USA P.J. Wakelyn 18.1 Introduction 18.2 Mattresses and/or Foundation (Box Springs)

266 270

xvii

Contents

18.2.1 US CPSC 18.2.2 CA BHFTI 18.3 Bedclothes 18.3.1 US CPSC 18.3.2 CA BHFTI 18.4 Upholstered Furniture 18.4.1 US CPSC 18.4.2 CA BHFTI (TB 117, TB 116) 18.4.3 Upholstered Furniture Action Council Voluntary Furniture Smoulder/ Cigarette Test 18.4.4 ASTM Standard 18.5 Children’s Sleepwear 18.5.1 1996 Amended Standard 18.5.2 New CPSC Data Collection Tool for Clothing-Related Burn Injuries to Children 18.6 Clothing Textiles 18.6.1 US CPSC General Apparel Standard (16 CFR 1610) 18.6.2 US CPSC Updated Standard (2008) 18.7 Carpets and Rugs 18.8 Ignition Sources 18.8.1 Cigarette Lighters 18.8.2 Candles and Candle Accessories 18.8.3 Matches 18.8.4 Cigarettes 18.9 Flame Retardant Chemicals 18.10 Summary and Conclusions References Chapter 19

270 271 272 272 272 275 275 280

282 283 283 283

284 284 285 285 286 287 287 287 287 287 288 289 290

Flame Retardancy of Cellulosic Fabrics: Interactions between Nitrogen Additives and Phosphorus-Containing Flame Retardants Sabyasachi Gaan, Gang Sun, Katherine Hutches and Mark Engelhard 19.1 19.2

Introduction Experimental 19.2.1 Material 19.2.2 Sample Preparation 19.2.3 Measurements 19.3 Results and Discussion 19.3.1. Flammability of Treated Fabrics 19.3.2. Surface Morphology of Char 19.3.3 ATR-FTIR Spectra of Char Surfaces

294 296 296 296 296 297 297 297 300

xviii

Contents

19.3.4 19.3.5

XPS Analysis of Char Mechanism of Formation of Surface Coating on the Char 19.4 Conclusion References Chapter 20

Chapter 21

Synergistic Flame Retardant Copolymeric Polyacrylonitrile Fibres Containing Dispersed Phyllosilicate Clays and Ammonium Polyphosphate A.R. Horrocks, J. Hicks, P.J. Davies, A. Alderson and J. Taylor 20.1 Introduction 20.2 Experimental Method and Results 20.2.1 Materials and Characterization 20.2.2 In-Situ Radical Polymerization of Nanocomposite Copolymers 20.2.3 Dope Blending of Clays 20.2.4 Polymer Spinning 20.2.5 Physical Characterization 20.5.6 Preparation and Flammability Testing of Flame Retarded Experimental Tows and Polymer Samples 20.6 Conclusions Acknowledgements References

303 305 305

307 308 308 310 313 313 314

320 328 329 329

Flame Retardance of Polyacrylonitriles Covalently Modified with Phosphorus- and Nitrogen-Containing Groups John R. Ebdon, Barry J. Hunt, Paul Joseph and Tara K. Wilkie 21.1 Introduction 21.2 Experimental 21.3 Characterization 21.4 Results and Discussion 21.5 Conclusions Acknowledgements References

Chapter 22

300

331 332 334 335 338 339 339

Novel Fire Retardant Backcoatings for Textiles M.A. Hassan 22.1 22.2

Introduction Experimental

341 342

xix

Contents

22.2.1 22.2.2 22.2.3 22.2.4 22.3 Result 22.3.1

22.3.2 22.3.3

22.3.4 22.3.5 22.3.6 References

Chapter 23

Materials Preparation of Flame Retardant Compounds Preparation of Coating Paste Characterization and Discussion Thermal Characterization of A1 and A5 Organophosphorus Compounds Pyrolysis Behaviour of Uncoated and Back-coated Cotton Samples Thermal Pyrolysis Process of Uncoated and Back-coated Polyacrylic Samples Flammability Properties Smoke Density Measurements Conclusion

342 342 345 346 346

346 348

351 354 355 357 358

Effect of Yarn, Fabric Construction and Colour in Respect of Red Reflectance and Pigmentation on the Thermal Properties and Limiting Oxygen Index of Flame Retardant Polypropylene Fabrics C. Kindness B.K. Kandola and A.R. Horrocks 23.1 23.2

Introduction Experimental Methods 23.2.1 Materials 23.2.2 Flammability Testing 23.2.3 Thermal Analysis 23.2.4 Air Permeability 23.2.5 Colour Measurement and Pigment Analysis 23.3 Results and Discussions 23.3.1 Commercial Fabrics 23.3.2 Experimental Fabrics 23.3.3 Air Permeability and the Effect on LOI 23.3.4 Effect of Colour on LOI 23.3.5 Thermal Analysis of Fabrics and Colour Pigments 23.4 Conclusions Acknowledgements References

359 362 362 365 366 366 366 367 367 368 369 370 372 373 376 376

xx

Contents

Fire Toxicity Chapter 24

Influence of Fire Retardants on Toxic and Environmental Hazards from Fires David Purser 24.1 24.2 24.3

Introduction Major Determinants of Toxic Product Yields Influence of Different Fire Retardant Systems on Toxic Product Yields and Toxic Hazards 24.3.1 Inert and Active Fillers 24.3.2 Phosphorus-based Systems 24.3.3 Nitrogen, Melamine and Melamine– chlorinated Phosphate Systems 24.3.4 Halogen Acid Vapour-phase Systems 24.3.5 Fluoropolymers and Ultrafine Particles (Nanoparticles) 24.3.6 Environmental Contamination by Dioxins and Furans from Halogenated Materials 24.4 Conclusions References Chapter 25

387 387 390 390 393 394

398 402 402

Assessment of Fire Toxicity from Polymer Nanocomposites Anna A. Stec and T. Richard Hull 25.1

Introduction 25.1.1 Fire Scenarios 25.1.2 The Steady-state Tube Furnace (ISO 19700) 25.1.3 Toxic Potency of Fire Effluent 25.2 Experimental 25.2.1 Materials 25.3 Results 25.3.1 Yields of Toxic Products from PA6 and PP 25.3.2 Fractional Effective Dose 25.3.3 LC50 of Different Polymeric Materials 25.4 Conclusions References

Subject Index

381 383

405 406 407 408 410 410 410 410 413 415 416 416 419

INTRODUCTION

Polymers and Fire T. RICHARD HULL AND ANNA A. STEC Centre for Fire and Hazards Science, School of Forensic and Investigative Sciences, University of Central Lancashire, PR1 2HE, Preston, Lancashire, United Kingdom

Unwanted fires account for significant losses to life and property. In the UK about 600 lives are lost each year,1 and the cost of unwanted fire has been estimated at about 1% of the UK’s gross domestic product.2 The vast majority of unwanted fires are fuelled by organic polymers and, as manufacturing technology has advanced, there has been a rapid shift from natural polymers (contained in wood, cotton, leather and wool) to synthetic polymers. Synthetic polymers are generally more flammable than their natural counterparts – polyethylene, polypropylene (PP) and polystyrene have calorific values comparable to that of petroleum. However, unlike natural polymers, which can only be fire retarded by coatings or other surface treatments, for most plastic materials the manufacturing process is ideal for the incorporation of fire retardants.

1 Hazards from Fire The flammability of a material is not an intrinsic property, like its density or heat capacity, but is dependent on the fire conditions. The apparent order of flammability of two materials may be reversed if tested under different conditions. Similarly, changing the material composition, for example by the addition of a fire retardant, will also change its reaction to fire behaviour. The incorporation of a nanofiller will reduce the dripping tendency. In one fire scenario, dripping away from a flame will reduce the ignitability, while in Fire Retardancy of Polymers: New Strategies and Mechanisms Edited by T Richard Hull and Baljinder K Kandola r Royal Society of Chemistry 2009 Published by the Royal Society of Chemistry, www.rsc.org

1

2

Introduction Heat

Fuel

Figure 1

Oxygen

The fire triangle.

another, drips, especially flaming drips, will cause downward flame spread, which significantly increases the fire hazard. The fire triangle (Figure 1) demonstrates the interdependence of the material properties with ventilation and heat. In general, fire growth is more favourable as the heat flux or oxygen supply increases, or if the material is more ‘‘flammable’’. However, excessive ventilation may remove heat from the flame, while additional heat may also result in melting or char formation, each of which could reduce fire growth. This scenario dependence ultimately favours certain materials under certain conditions. This is the heart of the difficulty in defining flammability, and explains why the materials development described in the different sections of this book often give apparently inconsistent results when tested under different conditions. Recent developments in flammability testing have brought us nearer to addressing the ultimate goal of predicting large-scale fire behaviour from smallscale tests, as have measurements of material properties coupled to models of full-scale fire behaviour. However, large-scale fires show considerable variation, so there is no universal benchmark against which to judge a material’s fire performance.

2 Fires and Fire Growth Fire tests that focus on particular fire stages should address the prevailing conditions appropriately. Most fires start from small beginnings. There may be an induction period (involving smouldering) before flaming ignition takes place, then a rise in temperature until ventilation-controlled burning occurs (usually 800–1000 1C), followed by decay as fuel is consumed, shown schematically in Figure 2. During the ignition phase, the impact of heat on a polymeric material causes an increase in temperature. If a sufficiently high temperature is reached then chemical bonds break and volatile fragments are produced. As they are hot, they are buoyant above the surface of the polymer. Once a sufficient concentration is reached, if the products are flammable then a flame may stabilize. Ignition may either be piloted by a flame or spark or be spontaneous; the latter typically occurs when the polymer surface is 200 1C hotter. There will be convected heat above the flame and radiative heat in all directions, including downwards, as shown in Figure 3.

Polymers and Fire

Figure 2

Stages in a fire.

Figure 3

Primary ignition process.

3

These heat-transfer processes are critical to the ignition and fire behaviour. Once ignited, the fire initially grows by a process of flame spread. The surfaces near the pyrolysis and flame zone are heated and decompose to form more flammable products. The flame spreads by piloted (i.e. flame) ignition of these areas beyond the burning zone. Hence flame spread is essentially a process of repeated ignitions, as depicted in Figure 4. Horizontal flame spread is relatively slow because the material ahead of the flame is heated by gas-phase conduction only, enhanced by downward

4

Introduction

Figure 4

Horizontal flame spread (slow).

Figure 5

Vertical flame spread (rapid).

radiation. Upward flame spread (Figure 5) is more rapid, because radiative, convective and some conductive heat transfer occurs. During the early growth phase, flaming is normally confined to the item first ignited. As flames become more than about 1 or 2 m high, radiative heat transfer to adjacent items becomes important, even to objects several metres away from the flame (Figure 6). These items decompose predominantly as a consequence of radiative heat transfer, pyrolyze and may then spontaneously ignite. As the flame reaches the ceiling, it spreads across it, dramatically increasing the radiant flux to other objects in the room. These then pyrolyze, filling the room with a flammable, or even explosive, fuel–air mixture. Once this occurs then the whole room will burn, and flashover is said to occur (Figure 7). At this point, the speed of flame spread will be faster than the running speed, and the fire can no longer be controlled. Burning will continue until the available fuel is consumed.

5

Polymers and Fire

Figure 6

Ignitability at a distance.

Radiation

Pyrolyzing fuel

Figure 7

2.1

Flashover conditions.

Conditions of Each Fire Stage

To simulate the effects of fire for materials development and testing, the test conditions should be related to the appropriate scenario.
Ignition. It is more difficult to obtain repeatable results from spontaneous ignition so, even though unwanted fires may result in this way, piloted ignition is generally the preferred scenario to assess the onset of flaming combustion. This is dependent on the ignition source (flame, cigarette, glow wire, etc.), sample size (1–10 cm) and ambient temperature.
Developing fire. The continuation of flaming combustion during fire growth involves an external heat flux of around 20–60 kW m2, which

6

Introduction

requires larger sample sizes (10 cm to 1 m), ambient temperatures above the ignition temperature (400–600 1C), with adequate ventilation.
Fully developed fire. The major stage of fire growth involves high external heat fluxes (450 kW m2), large sample sizes (1–5 m), ambient temperatures above the spontaneous ignition temperatures (4600 1C) and low ventilation. These conditions are not generally easy to replicate on a small scale, and materials which are required to perform well in developed fires normally need to be tested under these extreme conditions, and may perform differently in bench and large-scale scenarios.

2.2

Chemical and Physical Processes

The chemical composition of a polymeric fuel and the presence of fire retardants, additives, etc., are important in determining the degree to which flammable products will be released as the temperature increases. Untreated natural materials, such as wood, cotton and paper, tend to release flammable products and ignite at relatively low temperatures in comparison with synthetic materials [polyethylene, polyvinyl chloride (PVC), etc.]. However, the physical nature of the material also plays an important part (sometimes more so than the chemistry) in determining whether a material will reach decomposition temperatures. The thermal inertia (krc) is the product of the thermal conductivity, density and specific heat capacity. It dictates the time for the surface temperature to reach ignition temperature, describing the characteristics of materials according to their heat insulation or heat sink properties. A block of wood is more difficult to ignite with a small ignition source than are wood shavings. Cellular polymers of inherently combustible compositions (such as polyurethane foam) burn very rapidly in comparison to their solid counterparts because their heat insulation properties cause heat to be retained at the surface. The thermal inertia is low for insulating materials and high for heat-conducting materials. Ultimately, most fire science and hence most fire testing is focussed on specific protection goals, for good reasons. Common protection goals include preventing sustained ignition, limiting the contribution to fire propagation or acting as a fire barrier. Most of the better-established fire tests try to simulate a specific, realistic fire scenario and monitor a specific fire risk or hazard from a specific specimen within that scenario, rather than to determine the material’s properties. Furthermore, the way a specimen responds in a fire, or in a fire test, may make a significant contribution to the overall fire scenario. Hence, three general remarks can be made:
To compare the fire behaviour in different fire tests is difficult. Exact predictions often fail because different material properties determine the performance in different scenarios. However, rough correlations or correlations limited to specific classes of materials have been successful.
Scaling up and down is a key challenge in fire science, since the sample size plays such a major role. Typically, empirical approaches fail to predict fire

Polymers and Fire

7

behaviour satisfactorily, particularly attempts to span multiple orders of magnitude. Advanced predictive models have been developed which are moving towards reliable predictions of fire behaviour.
The interactions between properties of components and ‘‘intrinsic’’ material properties are complex and variable. Different polymers decompose in different ways and fire retardants act to inhibit the decomposition or flaming combustion processes. When a polymer is heated its chains will start to break down, which eventually results in the formation of volatile fuel molecules. The pyrolysis of a polymer, which turns polymer chains of 10 000–100 000 carbon atoms into species small enough to be volatilized, often involves breaking the polymer chain. In some cases, the chain releases groups from its ends most easily, known as end-chain scission or unzipping. In others, the chain breaks at random points along its length, known as random chain scission. A third process, in which groups that easily release are attached to the backbone as side chains, is known as chain stripping. This is often the preferred mechanism of the fire-retardant chemist, especially if the resulting chain may be prevented from undergoing chain scission to form volatiles or lose further substituents, and instead undergo carbonization that results in char formation. Thus, the conversion of organic polymer into volatile organic molecules may follow four general mechanisms. While some polymers fall exclusively into one category, others exhibit mixed behaviour. This process can be accelerated by chemical attack on the polymer chains, for example by atmospheric oxygen. In the presence of an ignition source, when the concentration of fuel molecules above the surface reaches a critical level, the proportion of their heat of combustion transferred back to the polymer is sufficient to replace the fuel by further pyrolysis. This is essentially the criterion for piloted ignition. It correlates well with the critical surface temperature for ignition. Once ignition has occurred, a proportion of the heat from the flame will be transferred back to an adjacent non-flaming part of the polymer surface, pyrolyzing the polymer and causing a repeat of the ignition process. This results in flame spread across the surface.

2.3

Studying Polymer Decomposition

Thermogravimetric analysis (TGA) provides a valuable insight into the decomposition behaviour of polymers under controlled conditions. The temperature at which significant mass loss occurs during decomposition in air gives an indication of the ignition temperature, as this is the point when a significant amount of fuel is lost from the polymer. This can be affected by gas-phase flame inhibitors and, to some extent, by the production of CO on the surface, which simultaneously reduces the oxygen and fuel concentrations. Once ignition has occurred, the mass loss in nitrogen is more representative of the fuel production rate, since the oxygen concentration under a flame is close to 0%.

8

Introduction

3 Fire Effluent Toxicity Analysis of fire statistics shows that most fire deaths are caused by inhalation of toxic gases.1 While some real-life fires may be represented by a single fire stage, most fires progress through several different stages.3 Burning behaviour and particularly toxic product yields depend most strongly on a few factors. Of these, material composition, temperature and oxygen concentration are normally the most important. The formation of carbon monoxide (CO), often considered the most toxicologically significant fire gas, is favoured by a range of conditions from smouldering to fully developed flaming. CO results from incomplete combustion, which can arise from:
Insufficient heat in the gas phase (e.g. during smouldering).
Quenching of the flame reactions (e.g. when halogens are present in the flame or excessive ventilation cools the flame).
The presence of stable molecules, such as aromatics, which survive longer in the flame zone and so give high CO yields in well-ventilated conditions, but lower than expected yields in underventilated conditions.4
Insufficient oxygen (e.g. in underventilated fires large radiant heat fluxes pyrolyze the fuel even though there is not enough oxygen to complete the reaction). The high yields of the asphyxiant gas CO from underventilated fires are held responsible for most of the deaths through inhalation of smoke and toxic gases,

Polymer Heat Pyrolysis

No

n-F

lam

ing

Ignition Flaming, Combustion (Heat flux 2-10x greater, Rapid increase in pyrolysis)

Op

en

Ve

ntil

atio

n

Restricted Ventilation Products rich in CO, smoke, organics and HCN if nitrogen present

Figure 8

Products rich in organics and partially oxygenated species

Effect of fire stage on toxic gas production.

Products mostly CO2 and H2O, (also SO2, NO2, acrolein and formaldehyde)

Polymers and Fire

9

but this underventilated burning is the most difficult to create on a bench scale. For most materials the yields of toxic species have been shown to depend critically on the fire conditions. Figure 8 illustrates the generalized change in toxic product yields during the growth of a fire from non-flaming through wellventilated flaming to restricted ventilation. Although the toxic product yields are often highest for non-flaming combustion, the rates of burning and the rate of fire growth are much slower, so underventilated flaming is generally considered the most toxic fire stage. Other toxic species include hydrogen cyanide (HCN, the other asphyxiant gas) and incapacitating irritants that cause blinding pain to the eyes and flooding of the lungs and respiratory tracts, which inhibit breathing and prevent escape. The wide variety of these irritants has led to groupings such as acid gases, organo-irritants and particulates, in order to estimate incapacitation.5 The effect of different fire conditions on the yields of these different toxicants is summarized in Figure 8. Data from large-scale fires6,7 show much higher levels of the two asphyxiant gases CO and HCN under conditions of reduced ventilation. It is therefore essential to the assessment of toxic hazard from fire that these different fire stages can be adequately replicated, and preferably the individual fire stages treated separately. Analysis of fire hazard requires data that describe the rate of burning of the material, and data that describe the toxic product yield of the material. This is best achieved using the steady-state tube furnace,8 in which the air supply and rate of burning are fixed as the sample is driven into a furnace, and then subjected to an increasing applied heat flux. Fire toxicity is also scenario dependent, but using this technique a clear relationship has been demonstrated between the yield of toxic products (for example in grams of toxicant per gram of polymer) and the fire condition for a given material composition.9 A more detailed account of current protocols in fire toxicity testing10 has recently been published.

4 Structural Deformation The increased use of polymer materials to replace structural members, such as the carbon fibre composites used in aircraft bodies, increases the importance of maintaining structural integrity during a fire. In many other cases, such as electrotechnical products, the failure of plastic components early on in a fire could radically alter the course of a fire, with potentially devastating consequences. As new materials with greater rigidity and structural integrity are being developed, synthetic polymer composites are increasingly being used to replace metal components. Incorporation, for example, of a 1% loading of a nanofiller can have a very large, beneficial effect on these physical properties.

5 Fire Retardant Strategies These can be broadly separated into those that block the fire physically, and those that use alternative chemical reactions to stop the material from burning.

10

Introduction

They are outlined here to set the context for the specific approaches described in detail in the individual sections.

5.1

Physical Action

There are several ways in which the combustion process can be retarded by physical action:
By cooling. Endothermic reactions cool the material.
By forming a protective layer. Obstructing the flow of heat and oxygen to the polymer, and of fuel to the vapour phase.
By dilution. Release of water vapour or carbon dioxide (CO2) may dilute the radicals in the flame so it goes out. For example, the most widely used fire retardant, aluminium hydroxide [Al(OH)3], breaks down endothermically to form water vapour, which dilutes the radicals in the flame, while the residue of alumina (Al2O3) builds up to form a protective layer. Unfortunately, relatively large amounts may be needed to be effective (up to 70%) and the freshly formed Al2O3 can lead to afterglow.11 180200  C

2AlðOHÞ3 ðsÞ ! Al2 O3 ðsÞ þ 3H2 OðgÞ DH ¼ þ1:3 kJ g1

5.2

Chemical Action


Reaction in the gas phase. The radical reactions of the flame can be interrupted by a flame retardant. The radical concentration falls below a critical value, and the flame goes out. The processes that release heat are thus stopped, and the system cools down. However, interfering with the flame reactions often results in highly toxic and irritant partially burnt products, including CO, which generally increase the toxicity of the fire gases while reducing fire growth.
Reaction in the solid phase. The flame retardants work by breaking down the polymer so it melts like a liquid and flows away from the flame (just like trying to light candle wax without a wick). Although this allows materials to pass certain tests, sometimes fire safety is compromised by the production of flammable drops. Char formation. Better solid-phase flame retardants are those which cause a layer of carbonaceous char to form on the polymer surface. This can occur, for example, by the fire retardant removing the side chains and thus generating double bonds in the polymer. Ultimately, these form a carbonaceous layer by forming aromatic rings.

Polymers and Fire

11

Char formation usually reduces the formation of smoke and other products of incomplete combustion. Intumescence. The incorporation of blowing agents causes swelling behind the surface layer, and provides much better insulation under the protective barrier. The same technology is used for coatings for protecting wooden buildings and steel structures.

5.3

Polymer Nanocomposites

Polymeric materials that contain fillers with at least one dimension of only a few tens of nanometres have opened up an enormous range of possibilities in fire retardant research. Fillers may have dimensions that extend over four orders of magnitude, and their effects include:





reinforcing organic char as a barrier layer; providing a catalytic surface to promote char-forming reactions; enhancing the structural rigidity of the polymer; changing the melt–flow properties of the polymer close to its ignition temperature; and
providing intimate contact between a fire retardant and the host polymer. Initially, investigations involved polymer–clay nanocomposites, but more recently investigations have included the use of single and multi-walled carbon nanotubes,12 and other nanoscopic fillers with potential fire retardant properties, including layered double hydroxides12 (or hydrotalcites), Al(OH)313 and others.

5.3.1 Types of Nanofillers for Fire Retardancy Although the effect of polymer–clay nanocomposites on fire behaviour was first investigated over two decades ago, the wider study of the full range of sizes, morphologies, chemistries and surface treatments of polymer nanocomposites has only just begun.

5.3.2 Filler Morphology Traditionally, particles with a platy morphology, and especially montmorillonite, have been investigated, as it was generally assumed that these would most easily assist in the formation of a barrier layer. The influence of nanocomposite formation and the different mechanisms of breakdown of different polymers make generalizations regarding filler morphology difficult. In some cases fillers with aspect ratios greater than 1000 have been successfully deployed to enhance fire retardancy. For example, a clay filler with a mean diameter of 25 mm has been used commercially to reduce the flammability of cable sheathing materials.14

12

Introduction

5.3.3 Filler Coating and Dispersion To produce appropriately dispersed polymer nanocomposites, it is generally necessary to add a compatibilizing agent, such as a surfactant, to the polar filler surface in order to insert it between the polymer chains. In partially ionic polymers, such as nylon, dispersion is much easier than in hydrophobic, crystalline polymers, such as isotactic PP. In these cases it is generally necessary to attach a grafting agent, such as maleic anhydride, onto the polymer to ensure adequate dispersion. While the mechanical properties of the polymer depend on adequate dispersion at ambient temperatures, the fire behaviour is a function of the dispersion of the nanofiller in the molten bubbling polymer. In many cases, the surfactant decomposes to leave the polar nanofiller. There has been intense speculation as to whether this results in incompatibilization and migration of the filler to the surface, or whether the preferential loss of the first few hundred nanometres of polymer results in accumulation of filler at the surface. In some cases no fire retardant effect is observed without adequate dispersion. In others it is evident, which suggests that dispersion occurs in the molten, decomposing bubbling polymer.

5.3.4 Effects of Nanofiller Composition on Thermal Decomposition Burning Behaviour The evolution products from nanocomposites made from polyethylene, ethylene vinyl acetate (EVA) and polystyrene, with organically modified clays, single and multi-walled carbon nanotubes and layered double hydroxides has been studied.12 It was found that the relative amounts and the identity of the degradation products change when both well-dispersed cationic and anionic clays are used, but there is no difference in the degradation products when carbon nanotubes are utilized. When the nanodimensional material is not well dispersed, the degradation products are not changed. Unlike clays, polymer-layered double hydroxide nanocomposites give reasonably good reductions in peak heat-release rate, even when nanodispersion has not been obtained. These data suggest that the enhancement in the fire behaviour must be, at least in part, due to different mechanisms for montmorillonite, layered double hydroxides and carbon nanotube-based nanocomposites.

6 Conclusions Fire is a complex process and no two real fires are identical. In developing materials with enhanced fire safety, such as lower ignitability, lower heatrelease rates during burning and lower fire toxicity, it is essential to relate the desired properties to the end-use scenario. It is normally the role of regulators to select appropriate test methods to protect people and property from the most likely fire scenarios.

Polymers and Fire

13

The thermal decomposition of polymers is a complex process, which may follow a number of different routes, depending on the material and the conditions. Polymers burn by breakdown of their long-chain structures, which releases fuel into the gas phase, where flaming combustion can occur. The mechanism of breakdown is often unique to a particular polymer and, in general, flame retardant methods cannot be directly exported from one polymer system to another. Fire retardants may be classified by their mode of action (physical or chemical, condensed phase or gas phase, char forming or intumescent, etc.) or by their physical or chemical structure. Polymer nanocomposites are an important new class of materials that offer thermal and mechanical properties not evident in their parent polymers, with great potential for fire retardancy. Some of their physical properties, for example the typically massive increase in viscosity modulus, may cause processing problems that prevent large-scale production using conventional extruders. The incorporation of additives, which have dimensions that range over four orders of magnitude, has a range of physical effects, such as barrier layer formation, loss of compatibilizer, migration to the surface, inhibition of bubble movement and reduction in the flow of the molten polymer. In addition, a number of chemical effects have been observed, including catalyzing decomposition reactions, promoting graphite formation and altering the decomposition pathway, which has been seen to influence the decomposition behaviour. The recent approach of preparing conventional fire retardants in nanoscopic form increases the range of chemical effects. The degree of dispersion has often been cited as a prerequisite for improved fire behaviour (typically a shorter time to ignition, but a lower peak of heat-release rate), but this is certainly not always the case. The controlling parameter is the degree of dispersion at the point of ignition, rather than that in the cold polymer, since either the compatibilizer may decompose, reducing the degree of dispersion, or the nanofiller may disperse under the more extreme agitation within the decomposing polymer. Measurements of rheological properties as a function of temperature have been shown to be an effective tool to demonstrate this.15 The complexities of fire behaviour and the difficulties in quantifying that behaviour in a scenario-independent way compound the problems of understanding the thermal decomposition of polymer nanocomposites. However, the large number of empirical studies that have produced encouraging results provides evidence that the future of fire retardancy will follow the nanocomposite route.9 Even if it is not yet possible to predict which type of nanofiller (in terms of chemistry, morphology, compatibilization and dimensions ranging over four orders of magnitude) and what degree of dispersion and filler loading are required for optimum performance. That nanofillers uniquely improve physical properties, while almost all other fire retardants worsen them, suggests that until such optimization has been reached, and the vast numbers of experiments required to achieve it have been undertaken, fire retardant formulations will be based on a combination of nanofiller and conventional flame retardant.

14

Introduction

References 1. Fire Statistics United Kingdom 2005, Office of the Deputy Prime Minister: London, April 2007. 2. The cost of fires – A review of the information available, Home Office Publications Unit, UK, 1997. 3. ISO TS 19706:2004 Guidelines for assessing the fire threat to people. 4. T.R. Hull, J.M. Carman and D.A. Purser, Prediction of CO evolution from small-scale polymer fires, Polym. Int., 2000, 49, 1259–1265. 5. ISO 13571:2007, Life-threatening components of fire – Guidelines for the estimation of time available for escape using fire data. 6. P. Blomqvist and A. Lonnermark, Characterization of the combustion products in largescale fire tests comparison of three experimental configurations, Fire Mater., 2001, 25, 71–81. 7. B. Andersson, F. Markert and G. Holmstedt, Combustion products generated by hetero-organic fuels on four different fire test scales, Fire Safety Journal, 2005, 40, 439–465. 8. T.R. Hull, J.M. Carman and D.A. Purser, Prediction of CO evolution from small-scale polymer fires, Polym. Int., 2000, 49, 1259. 9. A.A. Stec, T.R. Hull, K. Lebek, J.A. Purser and D.A. Purser, The effect of temperature and ventilation condition on the toxic product yields from burning polymers, Fire and Materials, 2008, 32, 49–60. 10. T.R. Hull and K.T. Paul, Bench-Scale Assessment of Combustion Toxicity – A Critical Analysis of Current Protocols, Fire Safety J., 2007, 42, 340–365. 11. T.R. Hull, R.E. Quinn, I.G. Areri and D.A. Purser, Combustion toxicity of fire retarded EVA, Polym. Degrad. Stab., 2002, 77, 235–242. 12. M.C. Costache, M.J. Heidecker, E. Manias, G. Camino, A. Frache, G. Beyer, R.K. Gupta and C.A. Wilkie, The influence of carbon nanotubes, organically modified montmorillonites and layered double hydroxides on the thermal degradation and fire retardancy of polyethylene ethylene-vinyl acetate copolymer and polystyrene, Polymer, 2007, 48, 6532–6545. 13. K. Daimatsu, H. Sugimoto, Y. Kato, E. Nakanishi, K. Inomata, Y. Amekawa and K. Takemura, Preparation and physical properties of flame retardant acrylic resin containing nano-sized aluminum hydroxide, Poly. Degrad. Stab., 2007, 92, 1433–1438. 14. T.R. Hull, D. Price, Y. Liu, C.L. Wills and J. Brady, An investigation into the decomposition and burning behaviour of ethylene-vinyl acetate copolymer nanocomposite materials, Polym. Degrad. Stab., 2003, 82, 365–371. 15. S. Nazare, T. R. Hull, B. Biswas, F. Samyn, S. Bourbigot, C. Jama, A. Castrovinci, A. Fina and G. Camino, Chapter 12, this book.

Fire Retardant Strategies

CHAPTER 1

Synergy between Nanometric Alumina and Organoclay in Conventional Fire Retardant Systems for Ethylene–Vinyl Acetate N. CINAUSERO,a J.-M. LOPEZ-CUESTA,a F. LAOUTID,a A. PIECHACZYKb AND E. LEROYa,c a

Ecole des Mines d’Ale`s, Centre des Mate´riaux de Grande Diffusion, 6, avenue de Clavie`res, 30319 Ale`s Cedex, France; b Nexans Research Center, 170 Av. Jean Jaure`s, 69633 Lyon Cedex 07, France; c Laboratoire de Ge´nie des Proce´de´s-Environnement-Agroalimentaire (GEPEA), 37, Bd de l’Universite´, 44606 Saint-Nazaire BP 420 Cedex, France

1.1 Introduction Hydrated mineral fillers like aluminium hydroxide (ATH) or magnesium hydroxide (MDH) are used in the cable industry as flame retardants for polyolefins such as ethylene–vinyl acetate (EVA) copolymers. The very high filler loadings usually required to obtain satisfactory fire properties,1,2 mean this results in a decrease in the mechanical performance of the materials. Nevertheless, enhancement of the efficiency of ATH or MDH may be achieved by partially substituting them with synergistic additives, in particular high-aspect ratio inorganic particles such as oMMTs or delaminated talcs.3,4 In addition to Fire Retardancy of Polymers: New Strategies and Mechanisms Edited by T Richard Hull and Baljinder K Kandola r Royal Society of Chemistry 2009 Published by the Royal Society of Chemistry, www.rsc.org

17

18

Chapter 1

improvement in the mechanical properties, the presence of such lamellar particles leads to an intumescence phenomenon that occurs before ignition in cone calorimeter tests. A foam-like charred structure is formed as a consequence of heterogeneous bubble nucleation, increased viscosity and the promotion of charring of the host polymer.4 This structure leads to the formation of a porous protective residue (mainly inorganic), which limits both heat transfer and the diffusion of fuel and oxygen.3 In a recent patent5 we showed that the addition of alumina nanoparticles improved the reactions to fire of flame retardant EVA compositions that contain metal hydroxide and oMMTs. In this chapter we present a detailed study of these complex systems and discuss the influence of the size of alumina particles.

1.2 Experimental 1.2.1 Materials EVA [Elvax 260, containing 28 weight percent (wt%) of vinyl acetate] was purchased from DuPont. Magnesium hydroxide (MDH; Magnifin H10, d50 ¼ 0.85 mm, specific surface area ¼ 10 m2 g
1) and oMMT (Nanofil 5: distearyldimethylammonium ion-exchanged bentonite) were supplied by Martinswerk (now Albemarle) and Su¨d Chemie (now Rockwood Holdings), respectively. Alumina particles of different physical properties were obtained from Degussa (ALU nano, Aeroxide Alu C, d50 ¼ 13 nm, SBET ¼ 86 m2g
1) and Alcan (ALU micro, d50 ¼ 0.47 mm, SBET ¼ 6.5 m2g
1), respectively.

1.2.2 Processing Blending of molten EVA copolymer with the different minerals was performed using a Haake internal mixer at 160 1C and 60 revolutions per minute (rpm) for 10 minutes. Thick (4 mm) sheets were then compression moulded at 160 1C at a pressure of 100 bars (1  107 Pa) for five minutes. These sheets were cut to the size required for the experiment to be performed. For all the different compositions studied, the total filler content was 10% or 60% w/w. As an example (EVA 40/ MDH 50/ALU nano 5/oMMT 5) means a formulation that contains 40% w/w of EVA, 50% of MDH, 5% of ALU nano and 5% of oMMT.

1.2.3 Testing E´piradiateur tests (AFNOR NF P 92-505) were carried out on 70  70  4 mm3 samples to determine the flammability and the self-extinguishability of the different compositions. The heat flux of e´piradiateur measured using a flux meter is around 30 kW/m2. This test allows the time to ignition (TTI) for the sample placed under the radiator (500 W) to be determined. After first ignition, the radiator is successively removed and replaced as soon as extinction occurs, the procedure being repeated for a period of 5 minutes. The mean inflammation

Synergy between Nanometric Alumina and Organoclay

19

period (MIP) and the number of ignitions (N) are then calculated from results obtained from four experiments for each formulation. Cone calorimeter tests (ISO 5660) were performed on filled polymer samples (100  100  3 mm3) placed horizontally, using a FTT cone calorimeter. Irradiances of 30, 50 and 70 kW/m2 were used. Ignition is piloted by a spark generator in contrast to the e´piradiateur test, during which ignition is spontaneous. TTI and peak of heat-release rate (PHRR) values are discussed later. The results given correspond to mean values obtained from two experiments for each formulation. Thermogravimetric analysis (TGA – Perkin Elmer PYRIS 1) was used to study the thermo-oxidative degradation of composites. Samples of typically 15 mg were placed in alumina crucibles and subjected to a temperature ramp from 25 1C to 700 1C in air at a heating rate of 5 1C min
1.

1.3 Results and Discussion To study the interactions between the various components of the formulations above, we performed three series of TGA. Figure 1.1 shows the effect of the presence of oMMT, ALU micro and ALU nano on the thermo-oxidative degradation of EVA. Pure EVA shows two main mass losses, the first one corresponding to EVA deacylation and the second one to EVA main-chain degradation. In the presence of oMMT, the first mass loss takes place at lower temperatures, while the second one is shifted towards higher temperatures, in agreement with literature results.6 The main explanation of the acceleration of the acetic acid loss is the catalytic effect of hydroxyl groups on the clay.6,7 In contrast, the presence of ALU nano does not affect the temperature of either of the two mass losses, while the presence of ALU micro leads to a slight change in the temperature of the second mass loss. Figure 1.2 displays the effect of nanofillers on the dehydration of MDH. Pure MDH dehydration takes place between 270 and 390 1C, with a maximum mass loss rate at 370 1C. In comparison, pure oMMT shows a mass loss between 200 and 390 1C, which corresponds to the degradation of its organic part, while ALU nano (and ALU micro) does not show any mass loss in this temperature range. When MDH is mixed with ALU nano (50/50 w/w powder mix), the dehydration of MDH starts at lower temperature, with a maximum mass loss rate at 355 1C. In comparison, when MDH is mixed with oMMT (50/50 w/w powder mix), the opposite occurs, with a maximum mass loss rate at 380 1C. In the meantime, the degradation of the organic part of oMMT does not appear to be influenced by the presence of MDH. Figure 1.3 shows the TGA and differential thermogravimetric (DTG) curves for the flame retardant formulations. The first mass loss, which corresponds to both the deacylation of EVA and the dehydration of MDH, is slightly influenced by the presence of ALU nano (the maximum mass loss rate is shifted towards lower temperatures). This may confirm its influence on the decrease of the MDH dehydration temperature observed in Figure 1.2. When both ALU

20

Chapter 1 100 90

Residual mass (%)

80 70 60 50 EVA 100

40

EVA 90 oMMT 10 30

EVA 90 ALU micro 10

20

EVA 90 ALU nano 10 EVA 90 ALU micro 5 oMMT 5

10

EVA 90 ALU nano 5 oMMT 5 0 250

300

350

400 Temperature (°C)

450

500

550

450

500

550

0

Mass loss rate (%/min)

-2 -4 -6 -8 EVA 100 -10

EVA 90 oMMT 10 EVA 90 ALU micro 10

-12

EVA 90 ALU nano 10 -14

EVA 90 ALU micro 5 oMMT 5 EVA 90 ALU nano 5oMMT 5

-16 250

300

350

400 Temperature (°C)

Figure 1.1

TGA and DTG curves for samples that do not contain MDH.

nano (or micro) and oMMT are present, a shoulder appears on the low temperature side of the mass loss rate peak. In this case, two phenomena can accelerate the mass loss: ALU nano seems to sharpen the reaction of deacylation of EVA catalyzed by oMMT, as observed in Figure 1.1; in addition, when incorporating both nanofillers, water release may be restricted, which contributes to the acceleration of mass loss at lower temperatures.

21

Synergy between Nanometric Alumina and Organoclay 100

Residual mass (%/min)

95

90

85 oMMT 80

ALU nano MDH

75

MDH - ALU nano MDH - oMMT

70 100

150

200

250 300 Temperature (°C)

350

400

450

350

400

450

0

Mass loss rate (%/min)

-0,5 -1 -1,5 -2

oMMT ALU nano

-2,5

MDH MDH - ALU nano

-3

MDH - oMMT -3,5 100

150

200

250

300

Temperature (°C)

Figure 1.2

Thermogravimetric effect of the presence of nanofillers on MDH dehydration.

Furthermore, the presence of oMMT shifts the first maximum mass loss rate towards a higher temperature, in agreement with the results of Figure 1.2 (dehydration of MDH at higher temperatures). The second mass loss peak is clearly shifted towards higher temperatures in the presence of oMMT, a shift due to EVA charring. From these TGA results we can make the assumption that the presence of ALU nano (or micro) in flame retardant formulations does not affect the flame

22

Chapter 1 100

Residual mass (%)

90 80 70 EVA 40 MDH 60 60

EVA 40 MDH 55 oMMT 5 EVA 40 MDH 50 ALU nano 10

50

EVA 40 MDH 50 ALU micro 5 oMMT 5 EVA 40 MDH 50 ALU nano 5 oMMT 5

40 250

300

350

400 Temperature (°C)

450

500

550

450

500

550

0

Mass loss rate (%/min)

-1 -2 -3 EVA 40 MDH 60 -4

EVA 40 MDH 55 oMMT 5 EVA 40 MDH 50 ALU nano 10

-5

EVA 40 MDH 50 ALU micro 5 oMMT 5 EVA 40 MDH 50 ALU nano 5 oMMT 5

-6 250

300

350

400 Temperature (°C)

Figure 1.3

TGA and DTG curves for samples containing MDH.

retardant action of oMMT, but can modify the dehydration of MDH, which starts at a lower temperature. Besides, oMMT used with alumina particles accentuates the deacylation of EVA in the low temperature range (250–340 1C), which could have a favourable effect on further charring due to double bond formation at lower temperature. The data obtained from the e´piradiateur tests are presented in Figure 1.4. The reference flame retardant formulation (EVA 40/MDH 60 w/w%) has the lowest TTI, which means the highest flammability, and the highest MIP, which means a poor auto-extinguishing ability. The introduction of oMMT clearly improves these two characteristics, while in the case of ALU (nano or micro) no significant improvement is observed. When both oMMT and ALU (nano or

23

Synergy between Nanometric Alumina and Organoclay 180

9,5 TTI épiradiateur (s)

170

9

MIP (sv)

160

8,5

TT I ( s)

150 8 140 7,5 130 7

120

6,5

110 100

6 EVA 40 MDH 60

Figure 1.4

EVA 40 MDH 55 oMMT 5

EVA 40 MDH 50 oMMT 10

EVA 40 MDH 50 Alu 10

EVA 40 EVA 40 MDH 50 MDH 50 oMMT 5 oMMT 5 Alu nano 5 Alu micro 5

E´piradiateur test results.

HRR (kW/m²)

500 450

30 kW/m²

400

50 kW/m²

350

70 kW/m²

300 250 200 150 100 50 0 0

Figure 1.5

200

400 Time (s)

600

800

Cone calorimeter HRR curves for the EVA 40/MDH 60 formulation.

micro) are present, the TTI increases significantly, compared to formulations that contain only oMMT. The best increase in TTI is obtained for ALU nano. The size of alumina particles is therefore an important parameter. Figures 1.5 and 1.6 show the cone calorimeter HRR curves obtained at various incident heat fluxes for (EVA 40/MDH 60), and (EVA 40/MDH

24

Chapter 1 500 30 kW/m²

HRR (kW/m²)

450 400

50 kW/m²

350

70 kW/m²

300 250 200 150 100 50 0 0

Figure 1.6

200

400 Time (s)

600

800

Cone calorimeter HRR curves for the EVA 40/MDH 50/oMMT 10 formulation.

50/oMMT 10), formulations, respectively. The behaviour shown in Figure 1.5 is ‘‘classical’’: when the incident heat flux decreases, the TTI increases and the PHRR decreases. In contrast, the presence of oMMT (Figure 1.6) leads to an unusual behaviour at low irradiance (30 kW/m2): the TTI becomes extremely long while the PHRR strongly increases. Similar behaviour was observed for the (EVA 40/MDH 55/oMMT 5) formulation, as shown in Table 1.1. This strong increase of PHRR at low incident-heat flux caused by the presence of oMMT is likely to be a problem for cable applications of the material. Effectively, studies8,9 have shown a good correlation between the PHHR in the cone calorimeter at low incident-heat flux (typically below 50 kW/m2) and the passing of the FIPEC cable test,10 which is a vertical tray test using a 20 kW burner. As Table 1.1 shows, the introduction of ALU nano allows the negative effect of oMMT on the PHRR at 30 kW/m2 to be decreased. In contrast, when ALU micro is used, the PHRR is increased, showing again that the size of alumina is an important parameter. As regards ignition, when incorporating nanoclays in flame retardant EVA, an increase of TTI occurs that we have already observed in a previous study.4 If we now focus on the TTI values obtained in the cone calorimeter test at 30 kW/m2, it is striking that they strongly differ from those obtained in e´piradiateur tests, although the incident heat flux is nearly the same. Contrary to what is observed for e´piradiateur tests (Figure 1.4), the introduction of ALU nano leads to a relative decrease of the TTI in cone calorimeter tests relative to that of the EVA/MDH/oMMT composition. These contradictory evaluations of the TTI are undoubtedly related to different experimental conditions: in the

Cone calorimeter results. 30 kW/m2

EVA EVA EVA EVA EVA EVA

40/MDH 40/MDH 40/MDH 40/MDH 40/MDH 40/MDH

60 55/oMMT 5 50/oMMT 10 50/ALU nano 10 50/oMMT 5/ALU nano 5 50/oMMT 5/ALU micro 5

50 kW/m2

70 kW/m2

TTI (s)

PHRR (kW/m2) TTI (s)

PHRR (kW/m2) TTI (s)

PHRR (kW/m2)

171 314 415 173 279 384

204 388 496 194 306 410

293 255 286 253 249 250

398 297 315 332 286 289

(+84%) (+143%) (+1%) (+63%) (+124%)

(+90%) (+143%) (–5%) (+51%) (+100%)

77 91 75 80 84 92

(+18%) (– 1%) (+4%) (+9%) (+19%)

(– 13%) (– 2%) (– 14%) (– 15%) (– 15%)

45 53 43 48 48 52

(+18%) (– 4%) (+7%) (+7%) (+17%)

(– 25%) (– 21%) (– 17%) (– 28%) (– 27%)

Synergy between Nanometric Alumina and Organoclay

Table 1.1

25

26

Chapter 1

case of e´piradiateur tests, the gases emitted from the sample are not aspirated, as in the cone calorimeter tests. To sum up, both the acetic acid and water are produced more efficiently owing to the effect of ALU on the deacylation of EVA catalyzed by oMMT observed in TGA, as well as the regulated water release from the dehydration of MDH. Therefore, we can assume that acetic acid and water may dilute the combustible gases at the surface of the sample during the pre-ignition period of e´piradiateur tests. This is likely to delay ignition, which in addition is not promoted by a spark in this test. Then the effectiveness with which ALU nano increases the TTI of a cable in a real fire will depend on the fire scenario. Nevertheless, the ‘‘static’’ conditions of the e´piradiateur are more likely than the forced flow of the cone calorimeter. Let us now come back to the PHRR values at 30 kW/m2. Comparison of the shape of the curves in Figure 1.6 suggests that, for this low external heat flux, a more important flux of combustible gas evolves at the time of ignition, and results in a strong PHRR. Such a ‘‘critical phenomenon’’ observed for (EVA 40/ MDH 50/oMMT 10) could be explained by a stronger migration of clay platelets11 towards the surface at low irradiance. This forms a protective layer before ignition that becomes ‘‘unstable’’ after ignition because of the additional external heat flux provided by the flame. Such a ‘‘critical phenomenon’’ is not observed for (EVA 40/MDH 60; Figure 1.5) and (EVA 40/MDH 50/ALU nano 10; Table 1.1) formulations, and is significantly reduced when both ALU and oMMT are present (EVA 40/MDH 50/ALU nano 5/oMMT 5; Table 1.1). Eventually, this last formulation is the best compromise concerning TTI and PHRR. In contrast, when ALU micro is used (Table 1.1), the PHRR is increased compared to that of the (EVA 40/MDH 55/oMMT 5) formulation, which confirms the size dependence of alumina particles on flammability properties.

1.4 Conclusion The effect of alumina particles on the thermo-oxidative degradation and the reaction to fire of conventional flame retardant formulations for EVA that contains MDH and oMMT has been studied. The introduction of alumina particles did not have any direct effect on the thermo-oxidative degradation of the EVA copolymer, but was shown to shift the dehydration of MDH towards lower temperatures. Besides, when mixed with oMMT, alumina particles may accentuate the deacylation of EVA catalyzed by oMMT. It was suggested that the restriction of water release and the acceleration of acetic acid loss could have an effect on the reaction to fire. Effectively, the e´piradiateur test showed a strong increase in the TTI in the presence of both alumina particles and oMMT, the best improvement being obtained for nano alumina particles. In addition, the use of alumina nanoparticles allowed the PHRR of EVA/MDH/ oMMT formulations to decrease at low external heat flux in cone calorimeter, which showed high values. This phenomenon is particularly relevant since the PHRR at low external heat flux is known to correlate with larger scale cable fire tests.

Synergy between Nanometric Alumina and Organoclay

27

References 1. F. Montezin, J.M. Lopez-Cuesta, A. Crespy and P. Georlette, Fire Mater., 1997, 21, 245–252. 2. R.N. Rothon, Particulate-Filled Polymer Composites, Longman Scientific & Technical, Harlow, Essex, England, 1995. 3. L. Ferry, P. Gaudon, E. Leroy, J.M. Lopez-Cuesta in Fire Retardancy of Polymers: New Applications of Mineral Fillers, ed. M. Le Bras, C.A. Wilkie, S. Bourbigot, S. Duquesne and C. Jama, The Royal Society of Chemistry, Cambridge, 2005, ch. 22 pp. 345–358. 4. L. Clerc, L. Ferry, E. Leroy and J.M. Lopez-Cuesta, Polym. Degrad. Stab., 2005, 88, 504–511. 5. A. Piechaczyk, J. Fournier, E. Tavard, J.M. Lopez-Cuesta, F. Laoutid and E. Leroy, European Pat., 2007, 1752490. 6. M. Zanetti, G. Camino, R. Thomann and R. Mu¨lhaupt, Polymer, 2001, 42, 4501–4507. 7. M.C. Costache, D.D. Jiang and C.A. Wilkie, Polymer, 2005, 46, 6947–6958. 8. M.M. Hirschler, Fire Technol., 1997, 33, 291–315. 9. M.M. Hirschler, Proceedings of Interflam’2001, Edinburgh, UK, 2001, Interscience Communications, London, UK, pp. 137–148. 10. S.J. Grayson, P. Van Hees, A.M. Green, H. Breulet and U. Vercellotti, Fire Mater., 2001, 25, 49–60. 11. M. Lewin, Polym. Degrad. Stab., 2006, 17, 758–763.

CHAPTER 2

Strained Organophosphorus Compounds as Reactive Flame Retardants for Polymeric Materials BOB A. HOWELL Center for Applications in Polymer Science, Central Michigan University, Mt. Pleasant, MI, USA

2.1 Introduction The development of new, more effective and environmentally friendly flame retardant strategies for styrenics is of increasing urgency as the criticism of the use of organohalogens, primarily brominated aromatics, becomes more prominent around the world.1–5 Flame retarding species that can be chemically incorporated into the polymer are particularly attractive, since volatility, blooming and loss during processing are not limitation issues for the use of the retardant. A goal of this work was to develop an organophosphorus (phosphate or phosphonate) that contains a thermally labile carbon–carbon bond that might be cleaved homolytically to generate a diradical that could act as an initiator for styrene polymerization.6,7 The use of such a compound as initiator would generate a polymer with a phosphorus-containing unit incorporated into the main chain. The polymer should display desirable flammability properties and be readily processable using standard techniques. It is possible that such a molecule with a strained carbon–carbon bond might also be reactive towards propagating radicals, i.e., it might act as a monomer. In this case, several phosphorus-containing units may Fire Retardancy of Polymers: New Strategies and Mechanisms Edited by T Richard Hull and Baljinder K Kandola r Royal Society of Chemistry 2009 Published by the Royal Society of Chemistry, www.rsc.org

28

Strained Organophosphorus Compounds as Reactive Flame Retardants

29

be incorporated into the polymer main chain. In fact, if the compound were to be sufficiently reactive as a monomer, the level of incorporation could be controlled by simply adjusting the concentration of the phosphorus compound in the polymerization mixture.

2.2 Experimental 2.2.1 Materials Benzpinacol was obtained by photoreduction of benzophenone. Benzophenone, dichloro(phenyl)phosphine and triethylamine were obtained from the Aldrich Chemical Company and used as received.

2.2.2 Initiator, 2,4,4,5,5-Pentaphenyl-1,3,2-Dioxaphospholane A solution of 1.00 g (2.71 mmol) of benzpinacol and 0.553 g (5.50 mmol) of triethylamine in 50 ml of anhydrous diethyl ether was placed into a 100 ml, threenecked, round-bottomed flask fitted with a magnetic stir bar, a pressure-equalizing dropping funnel and an Allihn condenser bearing a gas-inlet tube. Dichloro(phenyl)phosphine (0.49 g, 2.71 mmol) was added dropwise over a period of 0.25 hour. The mixture was stirred at solvent reflux for two hours, allowed to cool to room temperature and triethylammonium chloride was removed by filtration. The filtrate was subjected to rotary evaporation at reduced pressure to remove the solvent. The solid residue was crystallized twice from hexane to afford 2,4,4,5-pentaphenyl-1,3,2-dioxaphospholane (0.90 g, 70.5 % yield) as a white crystalline solid, Melting point (m.p.) 72 1C [differential scanning calorimetry (DSC)]; 1H nuclear magnitic resonance (NMR) (d, CDCl3), 7.01–7.42 (m, 25 H, aromatic protons); P-31 NMR (d, CDCl3), 22.1(s); Fourier transform infrared (FTIR) spectroscopy (cm 1, NaCl), 3028(vs) (aromatic C–H stretch), 2925[very strong (vs)] (aliphatic C–H stretch), 1601(vs) (aromatic nucleus), 1059 and 759 (O–P–O stretch and bending); mass spectrum m/z, (% of base), 77 (70), 106 (62), 141 (21), 182 (100), 473 (53).

2.2.3 Polymers 2.2.3.1

Polymerization of Styrene (1% 2,4,4,5,5-Pentaphenyl1,3,2-Dioxaphospholane as Initiator) at 70 1C

A polymerization tube fitted with a two-way stopcock was charged with inhibitor-free dry styrene monomer (10.0 g, 96.0 mmol) and 2,4,4,5,5-pentaphenyl1,3,2-dioxaphospholane (0.10 g, 0.21 mmol). The solution was degassed with pure dry nitrogen for 30 minutes and the tube placed in a constant-temperature

30

Chapter 2

oil bath at 70 1C for 15 hours. Size exclusion chromatographic (SEC) analysis of the polymer after purification by repeated dissolution in toluene followed by precipitation with methanol gave molecular weight (MW) 3.65  105, mean numerical molecular weight (Mn) 1.80  105 with a polydispersity of 2.03 and spectral properties 1H NMR (d, CDCl3) 1.58 (3H, methylene and methine protons), 6.8–7.5 (aromatic protons); P-31 NMR (d, CDCl3) 31.4(s); FTIR (cm 1, NaCl), 3025(vs) (aromatic C–H stretch), 2922(vs) (aliphatic C–H stretch), 1601(vs) (aromatic nucleus), 1492, 1452 (characteristic bands for polystyrene), 1059 and 759 (O–P–O stretch).

2.2.3.2

Preparation of Other Polymers

Polymers that contain various levels of phospholane were prepared in an analogous manner. A reference polymer that contained no phospholane fragments was prepared similarly using azobisisobutyronitrile (AIBN) as an initiator.

2.3 Results and Discussion The synthesis of an appropriate phosphorus compound was based on an earlier observation that when 2,3-diphenyl-2,3-butanediol was treated with phosphoryl chloride at 56 1C, a cyclic phosphate, which could be isolated and characterized, was formed.8,9 This ester contains hydrogen atoms alpha to the ester functionality, so it is unstable at higher temperature in the presence of base and readily undergoes elimination of phosphate to form the corresponding diene. A similar ester in which elimination is prohibited may be prepared from 1,1,2,2-tetraphenylethanediol (benzpinacol). Several suitable phosphorus esters may be prepared, depending on the reagents used. For example, treatment of the diol with dichloro(phenyl)phosphine affords 2,4,4,5,5-pentaphenyl-1,3,2-dioxaphospholane Scheme 2.1. In this case, the phosphine served as both reactant and solvent. The cyclic compound 2,4,4,5,5-pentaphenyl-1,3,2-dioxaphospholane can be obtained in good yield as a white crystalline solid with a melting point of 72 1C (DSC). The proton NMR spectrum contains absorptions for aromatic protons and

Scheme 2.1

Synthesis of 2,4,4,5,5-pentaphenyl-1,3,2-dioxaphospholane.

Strained Organophosphorus Compounds as Reactive Flame Retardants

31

the phosphorus spectrum consists of a singlet at d 22.1 parts per million (ppm), relative to the absorption for 85% aqueous phosphoric acid as external reference. The infrared spectrum of this contains an aromatic absorption band at 1601 cm 1, as well as strong bands for O–P–O absorption at 1059 and 759 cm 1. As noted above, this five-membered ring structure contains a sterically strained carbon–carbon bond that might be expected to undergo thermally induced homolysis at modest temperatures. This may be easily demonstrated for 2,4,4,5,5-pentaphenyl-1,3,2-dioxaphospholane using thermogravimetry. The compound readily undergoes thermal decomposition with a maximum rate of degradation at 120 1C and a degradation onset temperature of 70 1C. Isothermal thermogravimetry was used to determine a half-life for decomposition at various temperatures. For example, the half-life for its decomposition at 72 1C is 10 hours. The 10 hour half-life temperature is a convenient way to classify polymerization initiators. The ease with which the central carbon– carbon bond of this bond undergoes homolysis means it should function as a polymerization initiator. This is illustrated in Scheme 2.2 for the polymerization of styrene. If the compound functions as initiator, each polymer chain generated should contain one phosphorus moiety. Alternatively, it is possible that the phosphorus compound could also function as a co-monomer in the polymerization, i.e., the propagating poly (styryl) radical could add to the strained carbon–carbon bond (see Scheme 2.3). A series of mixtures was prepared in standard polymerization tubes. The first tube contained pure styrene monomer and was used as a control. The remaining tubes contained solutions of 1, 5 and 10% by mass of 2,4,4,5,5-pentaphenyl-1,3,2-dioxaphospholane in styrene monomer, respectively. Tubes were placed in an oil bath maintained at 70 1C. The progress of polymerization was

P O

P O

P O •

Scheme 2.2

(m + n)

O •

O •

O •

H H2

H2

* C C mCH C

O

P

O

H2 H2 H C CH C C n*

Initiation of styrene polymerization with 2,4,4,5,5-pentaphenyl-1,3,2dioxaphospholane.

32

Chapter 2

.

P

H2C CH

O

O H2C CH

H2C CH

Scheme 2.3

C O

P

O

C O

P

O

..

H2 . C CH

Copolymerization of styrene and 2,4,4,5,5-pentaphenyl-1,3,2-dioxaphospholane.

Table 2.1

SEC Characterization of poly(styrene) produced using 2,4,4,5,5-pentaphenyl-1,3,2dioxaphospholane as initiator at 70 1C.

Initiator present (weight %)

MN

MW

Polydispersity

0 1 5 10

Na 180 000 120 000 90 000

Na 365 400 302 400 273 600

Na 2.03 2.52 3.04

a

N: no polymer formed.

followed by removing aliquots of the mixture as a function of time for viscosity measurement. The aliquots were diluted with benzene (1 g of mixture per 100 ml of benzene). The relative viscosity increased as a function of time, which reflected the extent of polymerization. SEC data for the polymers formed are contained in Table 2.1. From the chromatographic data, it is apparent that thermal decomposition of 2,4,4,5,5-pentaphenyl-1,3,2-dioxaphospholane is effective in initiating styrene polymerization. In fact, when no phospholane was present in the monomer, no polymer was produced at 70 1C. In the presence of phospholane, polymer is readily produced under the same conditions. Further, the molecular mass of the polymer formed decreases regularly, as expected, as the concentration of initiator (phospholane) is increased. Analysis of the of polymers by both infrared and phosphorus-31 NMR spectroscopy indicated that phosphorus units were incorporated into the polymer main chain at a much higher level than can be accounted for by initiation, i.e., the phosphorus moiety probably functions as both initiator and co-monomer.

33

Strained Organophosphorus Compounds as Reactive Flame Retardants

2.3.1 Thermal Properties of Styrene Polymers Containing Phosphorus Units The thermal stabilities of the styrene polymers generated using one, five and ten mass percent 2,4,4,5,5-pentaphenyl-1,3,2-dioxaphospholane as initiator were examined using thermogravimetry. The relevant decomposition data are given in Table 2.2. From these results, it would appear that the thermal stability of the polymers that contain a phospholane unit is similar to that of poly(styrene), i.e., incorporation of the phospholane into the polymer main chain does not diminish the thermal stability of poly(styrene).

2.3.2 Evaluation of Flammability The flammability of these polymers was evaluated using the UL 94 vertical burn test. Results are presented in Table 2.3. Table 2.2

Comparison of the extrapolated onset and maximum decomposition temperature of styrene polymers produced in the presence of one, five, and ten percent 2,4,4,5,5-pentaphenyl-1,3,2-dioxaphospholane at 70 1C.

Phospholane in polymerization mixture (weight %)

Extrapolated onset temperature for decomposition (1C)

Maximum decomposition rate temperature (1C)

0 1 5 10

418.1 424.2 429.3 433.2

438.9 443.0 448.3 454.1

Table 2.3

Flammability behaviour of styrene polymers generated by initiation with 2,4,4,5,5-pentaphenyl-1,3,2-dioxaphospholane at 70 1C.

Observation Total flaming combustion for each specimen Total flaming combustion for all five specimens of any set Flaming and glowing combustion for each specimen after second burner Cotton ignited by flaming drips from any specimen Glowing or flaming combustion of any specimen to holding clamp Classification

Level of phospholane in polymerization mixture (Weight %) 0 30s 250s

1 30s 250s

5 30s 250s

10 30s 80s

60s

60s

60s

60s

YES

YES

YES

NO

YES

YES

NO

NO

94V-1

o94V-1

34

Chapter 2

The results in Table 2.3 suggest that, at the level of incorporation, the phospholane imparts a modest flame retardancy to poly(styrene). It may be anticipated that similar results can be obtained from other common evaluation techniques, such as limiting oxygen index (LOI) or cone calorimetry.

2.4 Conclusions 2,4,4,5,5-Pentaphenyl-1,3,2-dioxaphospholane undergoes thermally stimulated homolysis of a carbon–carbon bond at 70 1C to generate a diradical which efficiently initiates styrene polymerization. Polymers produced from initiation in this manner contain phosphorus in the main chain. As assessed by thermogravimetry, these polymers display thermal stability comparable to that of conventional poly(styrene), i.e., incorporation of phosphorus moieties into the polymer main chain does not diminish the thermal stability of the polymer. The presence of phosphorus in the polymer leads to a decreased flammability.

References 1. P. Georlette, J. Simons and L. Costa, ‘‘Fire Retardancy of Polymeric Materials’’, Marcel Dekker, Inc, New York, NY, 2000, p. 245–284. 2. B.A. King, in Modern Styrenic Polymers: Polystyrenes and Styrenic Copolymers, Eds., J. Scheirs, D.B. Priddy, John Wiley and Sons Inc, New York, NY, 2003, p. 685–701. 3. M. Simonson, Polym. Mater. Sci. Eng., 2000, 83, 90. 4. E. Weil, ‘‘An Attempt at a Balanced View of The Halogen Controversy,’’ Proceedings, 10th International Conference on Recent Advances in Flame Retardancy of Polymeric Materials,’’ Business Communications Company, Norwalk, CT, 1999. 5. W. Weil, Polym. Degr. Stab., 1996, 54, 125. 6. B.A. Howell and J. Uzibor, J. Therm. Anal. Cal., 2006, 85, 45. 7. B.A. Howell and J. Uzibor, J. Vinyl Addit. Tehcnol., 2006, 12, 198. 8. B.A. Howell, Y. Cui and D.B. Priddy, Thermochim. Acta, 2003, 396, 1673. 9. Y. Cui, ‘‘An Assessment of the Impact of Head-to-Head Placement on the Thermal Stability of Poly(styrene)’’, M.S. Thesis, Central Michigan University, Mt. Pleasant, MI, 2001.

CHAPTER 3

Amorphous Silicon Dioxide as Additive to Improve the Fire Retardancy of Polyamides G. SCHMAUCKS, B. FRIEDE, H. SCHREINER AND J.O. ROSZINSKI Elkem AS, Materials, P.O. Box 8126, Kristiansand, NO-4675, Norway

3.1 Introduction Fumed, precipitated and other types of silica are considered reinforcing fillers for rubber and plastics,1 but silica is not a flame retardant filler because it does not actively react through the release of water or another mechanism in fires. However, it was found that the addition of various types of silica (silica gel, fumed silica and fused silica) can significantly reduce the heat release and mass loss rate.2 It was concluded that the mechanism is based on physical processes in the condensed phase, rather than on chemical reactions. A number of papers at the Fire Retardant Polymers Meeting (FRPM) in Berlin in 2005 dealt with the effect of different silicon sources as flame retardant additives in polymer systems. Conflicting results were presented. Duquesne et al.3 reported that the addition of silica to an intumescent system did not improve the effect and that the rate of expansion was even decreased. The use of fumed silica (specific surface area 380 m2 g 1) in a textile application in combination with other halogen-free materials also did not lead to an efficient flame retardant system.4 Schmaucks et al. presented a synthetic amorphous silicon dioxide (specific surface area of only 20 m2 g 1) that showed very positive results in cone calorimeter investigations of ethylene–vinyl acetate (EVA)- and polyamide-6 Fire Retardancy of Polymers: New Strategies and Mechanisms Edited by T Richard Hull and Baljinder K Kandola r Royal Society of Chemistry 2009 Published by the Royal Society of Chemistry, www.rsc.org

35

36

Chapter 3

(PA6)-based thermoplastic compounds, alone or in combination with glass fibres or aluminium trihydrate, respectively.5 Transmission electron microscopy (TEM) investigations prove that the silica particles are well dispersed, which seems to be essential to create a material that has excellent mechanical properties and high fire resistance at the same time. The evaluation of this material, SIDISTARs, was extended to other types of polyamides. The new results indicate that the efficiency of this silica as flame retardant additive depends on the polyamide type. Dispersion of the silica particles and the viscosity of the composite seem to play the most important roles. Fillers and additives used to improve the fire retardancy of polymeric materials act either by a solid-phase or a gas-phase mechanism. Some might even show both modes of action. Silica will stay intact and will not react in any fire scenario. It is not expected that it will show an effect other than the dilution of the polymer, which reduces the amount of fuel available for consumption during the fire. It might, however, in combination with other ingredients, contribute to char formation and therefore increase the residue amount and reduce the heatrelease rate (HRR) and the total heat released (THR). In a number of studies, which can be found in the literature, conflicting results have been reported.6 Primary particle size and dispersion are seen as the main reasons for the different behaviour of silica in these investigations, but to really understand these different effects it is necessary to look at the nature and morphology of the silica types used.7 A possible differentiation of silica types according to Chemical Abstract Service (CAS) Numbers is shown in Figure 3.1. Of course, as any list, this one is also disputable, but it at least shows the large number of different silica types available. As examples to illustrate the different morphologies, TEM images of precipitated silica (CAS no. 112926-00-8), pyrogenic or fumed silica (CAS no. 112945-52-5), with two different specific surface areas according to the Brunauer, Emmett and Teller (BET) method, and silica fume (SIDISTARs; CAS no. 69012-64-2), are shown in Figures 3.2–3.5. It is clear that the precipitated and the fumed silica consist of very small, irregularly formed primary particles, which are partly fused together and form agglomerates, or chains. They also have a narrow particle size distribution. The strong interactions between the small particles make it difficult, or even impossible, to disperse them as primary particles in a polymer matrix using conventional mixing technology. However, silica fume consists of spherical particles with a broad particle-size distribution, which minimises the particle interaction and allows easy dispersion down to primary particles in a polymer composite. In addition to the differences in morphology, we need to consider the chemical behaviour caused by the silanole groups at the silica surface. Precipitated silica has the largest number of silanole groups, making it very hydrophilic, hygroscopic and reactive towards silanes. Commercially available precipitated silica contains around 5% water. Fumed silica has about half the silanole group density than does precipitated silica, is still hydrophilic and contains about 2% water. The silica fume used in this investigation, however, although hydrophilic, is not hygroscopic and contains less than 0.8% water.

Synthetic Amorphous Silica (7631-86-9)

Wet

Thermal

Surface modified Silica (67762-90-7) (68611-44-9) (68909-20-6)

Amorphous Silica (7631-86-9)

Natural

Crystalline Silica

By-Products

Silica Gel (112926-00-8)

Pyrogenic Silica fumed silica (112945-52-5)

Diatomaceous earth (61790-53-2)

Fused Silica (60676-86-0)

Precipitated Silica (112926-00-8)

Electric-Arc Silica

Calcined (91053-39-3)

Silica Fume Microsilica (69012-64-2)

Plasma Silica

2) Flux-Calcined (68855-54-9)

Cristobalite (14464-46-1)

Quartz (14808-60-7)

Tridymite (15468-32-3)

Keatite (17679-64-0) Coesite(13778-38-6) Stishovite(13778-38-5) (Melanophlogite)

Figure 3.1

Amorphous Silicon Dioxide as Additive to Improve the Fire Retardancy

Silica CAS 7631-86-9

Classification of silica types according to CAS numbers. 37

38

Chapter 3

Figure 3.2

TEM image of precipitated silica (BET 190 m2 g 1) – 19.5 kx.

Figure 3.3

TEM image of fumed silica (BET 200 m2 g 1) – 19.5 kx.

Amorphous Silicon Dioxide as Additive to Improve the Fire Retardancy

Figure 3.4

TEM image of fumed silica (BET 50 m2 g 1) – 19.5 kx.

Figure 3.5

TEM image of silica fume (SIDISTARs; 20 m2 g 1) – 41.5 kx.

39

40

Chapter 3

3.2 Experimental 3.2.1 Materials The materials used are listed here and in Table 3.1. PA12 – Grilamids L20 G natural (EMS Chemie AG) PA6.6 – Durethans A30 (Lanxess AG) Glass fibres – EC10 (Glasseidenwerk Oschatz) Synthetic amorphous silica (silica fume) – SIDISTARs T120 (Elkem AS, Materials)  Surface-modified silica fume (with 3-aminopropyltriethoxysilane) – SIDISTARs T120XP.

   

3.2.2 Sample Preparation All compounds were prepared by extrusion in a co-rotating twin-screw extruder (L/D 40) with a screw diameter of 25 mm at a speed of 300 revolutions per minute (rpm) at 240–260 1C for PA6.6 and 200–220 1C for PA12, respectively, followed by granulation. The polymers were dried prior to compounding. The specimens for mechanical tests, as well as the plaques (100  100  4 mm) for cone calorimeter measurements, were prepared by injection moulding. All specimens were conditioned for 72 hours at standard climate 23/50 according to ISO 291, prior to the mechanical tests.

3.2.3 Test Methods Tensile testing was carried out according to DIN EN ISO 527 and Charpy impact test according to DIN EN ISO 179-1eA and DIN EN ISO 179-1eU, respectively. 20 mm specimens were used for the vertical burning test, according to UL94. Cone calorimeter measurements were performed according to ISO 5660. Table 3.1

List of compounds.

Materials

PA6.6 Glass fibres Silica fume Silica fume, surface modified PA12 Glass fibres Silica fume Silica fume, surface modified

Sample No. and composition in weight % (w/w) 1.1 100 0 0 0 2.1 100 0 0 0

1.2 90

1.3 80

1.4 70

10 0 0

20 0 0

30 0 0

2.2 90

2.3 80

2.4 70

10 0 0

20 0 0

30 0 0

6.1 95 0 5 0 7.1 95 0 5 0

6.2 90

6.3 80

0 10 0

0 20 0

7.2 90

7.3 80

0 10 0

0 20 0

6.4 95

6.5 90

3.1 80

3.2 80

0 0 10

10 10 0

0 20 0

4.1 80

4.2 80

4.3 90

4.4 80

10 10 0

0 20 0

10 0 10

0 0 5

5 0 5

3.3 90

3.4 80

3.5 50

10 0 10

30 20 0

4.5 50

5.1 90

5.2 80

30 20 0

0 0 10

0 0 20

5 0 5

41

Amorphous Silicon Dioxide as Additive to Improve the Fire Retardancy

3.3 Results and Discussion Results of the evaluation of silica fume addition (SIDISTARs) into PA6 and PA11, alone or in combination with varying amounts of glass fibres, have been reported previously.5 In the research reported herein, its application in PA6.6 and PA12 was investigated. In Table 3.2 the tensile modulus dependence on the silica concentration is summarised. As expected, the modulus increases with increasing amounts of the silica. The results of the impact test are given in Figures 3.6 and 3.7. Surprisingly, it was found that up to 20% of silica can be added to PA12 without any loss in impact strength, whereas an addition of only 5% to PA6.6 leads to a significant reduction in the toughness of the composite. Surface modification of the silica leads to an increase in the compound’s toughness. The fire retardancy of the composites was evaluated by cone calorimeter and the UL 94 burning test. The results of the UL 94 test are summarised in Table 3.3. In PA6.6 the addition of 20% silica leads to a reduction of the after-flame time from 140 to 96 seconds. For comparison, the addition of 20% glass fibres increases the after-flame time to 852 seconds. In PA12, 20% silica reduces the after-flame time from about 591 to 36 seconds, whereas 20% glass fibres reduce the after-flame time to 528 seconds. Cone calorimeter measurements were carried out at three different irradiations, 35, 50 and 70 kW m 2. Selected results of duplicate or triplicate measurements, respectively, are shown in Figures 3.8–3.12. Surprisingly, it was found that the effect of the silica seems to be different in PA6.6 and in PA12.

Table 3.2

Tensile modulus of silica filled polyamides. Tensile modulus according to DIN EN ISO 527-2 Et (Mpa) Standard deviation

Mixture

Composition (w/w)

Mean value

PA12 PA12/T120 PA12/T120 PA12/T120

95/ 90/10 80/20

1487 1503 1572 1792

18 13 37 70

PA66 PA66/T120 PA66/T120 PA66/T120

95/5 90/10 80/20

2942 3072 3414 3620

74 191 154 139

PA66/T120XP PA66/T120XP

95/5 90/10

3110 3389

84 82

42

Chapter 3

Charpy impact toughness acu / kJ/m2

400 350 300 250 200 150 100 50

PA12 + T120

Figure 3.6

PA66 + T120

90/10

95/5

100/0

--

80/20

90/10

95/5

100/0

--

80/20

90/10

95/5

100/0

0

PA 66 + T120XP

Charpy impact strength dependence on silica concentration.

notched impact toughness ack / kJ/m2

8 7 6 5 4 3 2 1

PA12 + T120

Figure 3.7

PA66 + T120

90/10

95/5

100/0

--

80/20

90/10

95/5

100/0

--

80/20

90/10

95/5

100/0

0

PA 66 + T120XP

Notched Charpy impact strength dependence on silica concentration.

43

Amorphous Silicon Dioxide as Additive to Improve the Fire Retardancy

Table 3.3

20 mm vertical burning test according to UL 94 V. Vertical burning behaviour according to UL 94

Polymer Sidistars T120

Sidistars T120XP

Total after-flame Materials classification timea [s]

Type

%

Glass fibre

1.1 1.2 1.3 1.4

PA66

100 90 80 70

0 10 20 30

0 0 0 0

0 0 0 0

140 66 852 775

–b V2 –b –b

2.1 2.2 2.3 2.4

PA12

100 90 80 70

0 10 20 30

0 0 0 0

0 0 0 0

591 501 528 240

–b –b –b –b

3.1 3.2 3.3 3.4 3.5

PA66

80 80 90 80 50

10 0 5 10 30

10 20 0 0 20

0 0 5 10 0

175 96 213 72 533

–b V2 –b V2 –b

4.1 4.2 4.3 4.4 4.5 5.1

PA12

80 80 90 80 50 90

10 0 5 10 30 0

10 20 0 0 20 0

0 0 5 10 0 10

724 36 87 300 62 72

–b V2 V2 –b V2 V2

Sample No.

a

Burning time after first and second ignition, five specimens. Outside classification range of UL 94.

b

In PA12 the silica behaves like an inert filler without any specific barrierforming properties. A rather high HRR indicates that the material will not be self-extinguishing in flammability testing. 20% silica in PA6.6 causes a reduction in HRR of 52% at 35 kW m 2 and 65% at 70 kW m 2. This significant increase in flame retardancy effect with increasing radiation is typical for materials that form surface barriers in a cone calorimeter. From the effects of 20% SIDISTAR T120 in PA12 and in PA6.6, it is obvious that the silica worked crucially better in PA6.6 than in PA12 (see Figure 3.12). In PA6.6 the better barrier was formed. Barrier formation is controlled by several processes during combustion, such as demixing, selforganisation, ablative re-assembly, decomposition, etc. It is assumed that the barrier formation depends on different parameters, such as the various repulsive and attractive interactions between the components, viscosity of the melt, pyrolysis temperature, bubbling due to decomposition, etc. Given these details, it is surprising that such a clear difference for SIDISTAR in these two polyamides occurred, since most of the properties for PA6.6 and PA12 are not very different. A different distribution of the silica particles in PA6.6 and in PA12 was seen as the most reasonable explanation for the observed different fire behaviours. This was investigated by TEM, and the results are shown in Figures 3.13 and 3.14.

44

Chapter 3 2500

150

2000

120

1500

90

1000

60

500

THR / MJ m-2

HRR / kW m-2

35

kW/m2

30 HRR THR

0

0 0

Figure 3.8

100

200

300 Time / s

400

600

500

HRR and THR for PA12 at 35 kW m 2.

150

2500

2000

120

1500

90

1000

60

30

500 HRR THR 0 0

Figure 3.9

100

200

300 Time / s

400

500

HRR and THR for PA12:silica fume 80:20 at 35 kW m 2.

0 600

THR / MJ m-2

HRR / kW m-2

35 kW/m2

45

Amorphous Silicon Dioxide as Additive to Improve the Fire Retardancy 2500

150

120

1500

90

1000

60

500

30 HRR THR

0 0

100

200

300 Time / s

0 600

500

400

HRR and THR for PA6.6 at 35 kW m 2.

2500

150 HRR

HRR / kW m-2

35 kW/m2

THR

2000

120

1500

90

1000

60

500

30

0 0

Figure 3.11

100

200

300 Time / s

400

500

0 600

HRR and THR for PA6.6:silica fume 80:20 at 35 kW m 2.

THR / MJ m-2

Figure 3.10

-2

2000

THR / MJ m

HRR / kW m-2

35 kW/m2

46

Chapter 3 Peak of heat release rate 3000 PA 12

PA 12/T 120 (80/20)

PA 66

PA 66/T 120 (80/20)

HRR / kW m-2

2500

2000

1500

1000

500

0 35

50

70

35

50

70

35

50

70

35

50

70

-2

External Heat Flux / kW m

Figure 3.12

HRR for PA12 and PA6.6 at three different irradiations.

Figure 3.13

TEM of PA6.6 containing 20% silica fume.

No differences in dispersion could be observed. The conclusion is therefore that reasons other than silica dispersion cause the differences in fire retardancy. When comparing the effect of 10% SIDISTAR T120 with 10% surfacemodified silica (SIDISTAR T120XP) in PA6.6, it was found that the

Amorphous Silicon Dioxide as Additive to Improve the Fire Retardancy

Figure 3.14

47

TEM of PA12 containing 20% silica fume.

unmodified silica has a more pronounced barrier effect. The peak HRR was reduced by one-third for T120, but only by one-quarter for T120XP. The surface-modified silica also showed almost no reduction in total heat evolved (THR) compared to T120. The decrease in peak HRR was accompanied by a corresponding decrease in maximum carbon monoxide (CO)- and smoke-release rate because of the reduced amount of polymer consumed by the fire.

3.4 Conclusion Synthetic amorphous silicon dioxide with a BET surface area of 20 m2 g 1 (silica fume – trade name SIDISTARs T120) shows different behaviour in terms of physical properties and fire retardancy in PA12 and PA6.6. In PA12 20% of the silica could be added without loss of impact strength, but it only shows the behaviour of an inert filler in the cone calorimeter test. In PA6.6, the addition of only 5% silica leads to a significant reduction in impact strength, but the composite possesses extraordinary flame retardancy when tested according to ISO 5660 in the cone calorimeter. The difference in behaviour cannot be explained by differences in silica dispersion.

Acknowledgement The authors thank Mrs Krajewsky and Mr Thieroff, Kunststoffzentrum, Leipzig, for the sample preparation and the mechanical and UL 94 tests, Dr Schartel of the Federal Institute for Materials Research and Testing (BAM), Berlin, for the cone calorimeter tests and Mr Seydewitz, University Halle, for the TEM investigations.

48

Chapter 3

References 1. Georg Wypych Handbook of Fillers, Plastics Design Library, Toronto New York 1999, p. 131 ff. 2. Kashiwagi et al. Journal Fire & Materials, 6, Nov./Dec. 2000, p. 277 ff. 3. Duquesne et al. Proceedings of FRPM’05, Berlin 2005, paper 2_O_1. 4. A.R. Horrocks et al. Proceedings of FRPM’05, Berlin 2005, paper 5_O_1. 5. G. Schmaucks et al. Proceedings of FRPM’05, Berlin 2005, paper 4_O_4. 6. T. Kashiwagi, A.B. Morgan, J.M. Antonucci, M.R. VanLandingham, R.H. Harris, W.H. Awad and J.R. Shields, J. Appl. Polym. Sci, 2003, 89(8), 2072–2078. 7. T. Kashiwagi, in Flame Retardant Polymer Nanocomposites, ed. A.B. Morgan and C.A. Wilkie, Wiley, Hoboken, 2007, p. 286–288.

CHAPTER 4

Use of Organosilicone Composites as Flame Retardant Additives and Coatings for Polypropylene + A. SZABO´, K. KISS AND G. MAROSI B.B. MAROSFOI, Budapest University of Technology and Economics, Department of Organic Chemical Technology, Budafoki u´t 8., H-1111 Budapest, Hungary

4.1 Introduction Organosilicone polymers and their organic–inorganic hybrid structures are promising components of flame retardant formulations. This is attributable to their superior properties, compared to polymers with an organic main chain, in terms of their thermal stability and electrical properties. Industrially, the most important silicone is polydimethylsiloxane (PDMS). Its conversion into polyborosiloxane (pBSil) via a polycondensation reaction has attracted interest in both the academic and industrial fields. The mechanical and rheological properties of pBSil have been widely studied, while their behaviour in thermooxidative atmosphere has been investigated less.1,2 These siloxane compounds may act during combustion as a shielding powder (PDMS) or as a protective ceramic layer (pBSil). The boron atom in pBSil is an essential component of ceramic layer formation, as without it the PDMS transforms into a fine silica powder.3 Advantageous synergetic actions between pBSil and ‘carbon-based’ intumescent flame retardants have also been reported – the pBSil acts as an additive that enhances the consistency of char.4 Fire Retardancy of Polymers: New Strategies and Mechanisms Edited by T Richard Hull and Baljinder K Kandola r Royal Society of Chemistry 2009 Published by the Royal Society of Chemistry, www.rsc.org

49

50

Chapter 4

Recent advances in flame retardant polymeric materials have focused on the flame retardancy of polymer–clay nanocomposites (PCNs). The reducing effects of clays in PCN on heat-release rate (HRR) are well known. Moreover, the clays in combination with intumescent flame retardants have a significant influence on the char structure through their bubble-nucleating effect.5 The needle-like sepiolite clay may promote the formation of increased amounts of char of better uniformity and strength. Different boron-containing compounds, such as metal borates (zinc, aluminoand magnesium borate) and organoborates [like melamine–borate (MB)], are increasingly used in polymers, particularly for fire retardant purposes. They may act as smoke suppressants, afterglow suppressants, corrosion inhibitors or synergistic agents.6–8 Inorganic boron salts act in the condense phase by changing the degradation pathway of polymers and thus promoting the formation of char and the reduction of carbon dioxide (CO2) and carbon monoxide (CO).9,10 Although all three types (polyorganosiloxanes, PCNs and MBs) have been used as synergists in fire retardant compositions, their combination without fire-retardant active atoms (e.g. phosphorous, bromine) has not been investigated yet. In this study of pBSil, various composites that contain needlelike clay and MB were prepared and characterized as flame retardant components in polypropylene (PP). The combustion characteristics of the relevant polymer blends and composites were compared to those of pristine PP. Contact angle measurements were made, as well as the fire retardant characterisation, to determine the multifunctional character of the developed pBSil composites.

4.2.

Experimental

4.2.1 Materials PP of Moplen HP400R type was received from Basell Polyolefins, MFI: 25 (230 1C/2,16 kg). Hydroxyl-terminated PDMS) {HO–[Si(CH3)2O)n–H]} with a viscosity range of 400–650 mPa
s was received from Wacker Silicone GmbH (Wacker Silicone OEL CT 601 M). Boric acid (BA), melamine and absolute ethanol were received from Sigma Aldrich and used without further purification. The MB was prepared in our laboratory and used as an incorporated additive. For the MB preparation, first the melamine (12.6 g, 0.1 mol) was dissolved in 600 ml boiling distilled water and the BA (12.2 g, 0.2 mol) in 100 ml distilled water. After mixing these, the solution was left to cool to room temperature under constant stirring. The white flocculated MB was completely precipitated then decanted and filtered. The obtained MB was then washed with cool water and left at room temperature to dry. The needle-like clay sepiolite in organomodified form (OSEP) was supplied by Tolsa (trade-mark: Pangel B40). a,a 0 -bis(t-butylperoxy)diisopropylbenzene (trade mark Luperox

51

Use of Organosilicone Composites as Flame Retardant Additives

F90P from Elf Atochem) was used as the peroxide radical initiator. When applied, the amount was 0.5% of the pBSil content of the composites.

4.2.2 Sample Preparation The formation of a boron linkage between the PDMS chain required a treatment in boiling ethanol. The BA was dissolved in boiling ethanol, and then the PDMS added dropwise. The mixture, under continuous stirring, was heated until a transparent solution was obtained. The reaction was carried out in excess BA. To apply additives and obtain a homogeneous mixture, MB and/or OSEP was added before the removal of ethanol. The majority of ethanol was removed by heating the samples up to 140 1C and any remaining ethanol subsequently removed by evaporation in a vacuum chamber at 60 1C. The pBSil-based samples were solidified after the evaporation. The compositions of prepared samples are summarised in the Table 4.1.

4.2.3 Preparation of PP Compounds MB and OSEP that contained polyborosiloxane (pBSil–MB–OSEP) systems were introduced into PP. The amount of (pBSil–MB–OSEP) system in PP was increased by a multistep process in the mixing chamber of Brabender Plastograph. A mixing temperature of 190 1C and a mixing speed of 60 min1 were used, the samples were homogenised for 10 minutes and then melt-compressed at 200 1C.

4.2.4 Preparation of Composites with Multilayer Structures The efficiency of the pBSil–OSEP–MB system when structured as multilayers was also investigated. The upper layer was a pBSil-based composite, while the internal layer was pristine PP. The multilayer structured samples were prepared in a Colin P200E type press at 195 1C and 5.25 kPa (50 bars) for 10 minutes by compression of a 10  10  0.4 cm PP sheet and a pBSil–OSEP–MB 10  10 sheet. The thickness of the pBSil-based outer layer was 0.5 mm.

Table 4.1

Composition of pBSil and its composites with MB and/or organomodified sepiolite (OSEP).

Sample code PDMS pBSil PBSil–MB(5%) PBSil–OSEP(5%) PBSil–MB(5%)–OSEP(5%)

PDMS (wt%) 100 90.9 86.4 86.4 81.8

BA (wt%)

MB (wt%)

OSEP (wt%)

– 9.1 8.6 8.6 8.2

– – 5 – 5

– –– 5 5

52

Chapter 4

4.2.5 Characterisation The thermal stability of materials was examined by thermogravimetric analysis (TGA: Setaram Labsys TG DTA/DSC) under air atmosphere and with a heating rate of 10 1C per minute. The combustion parameters [HRR, time-to-ignition (TTI) and total heat release (THR)] of samples were determined by a mass loss calorimeter (according to ISO 13927, FTT Inc.) under a heat flux of 50 kW m2. Contact angle values were determined goniometrically by the sessile drop method using 5 ml volume of water in a Kruss Contact Angle Meter.

4.3 Results and Discussion 4.3.1 Thermo-Oxidative Stability The analyses to determine the influence of boron inclusion on the thermooxidative stability of PDMS were designed considering former experiments to determine the role of different end groups on its thermal degradation.11,12 The TGA results in Figure 4.1 show, in agreement with the literature, that the oxidative thermal degradation of PDMS takes place in two stages. In the first

Figure 4.1

TGA and differential thermal analysis (DTA) curves of PDMS and pBSil polymers and systems that contain 5 wt% OSEP and/or MB component in air, with a heating rate of 10 1C per minute.

53

Use of Organosilicone Composites as Flame Retardant Additives

Table 4.2

Characteristic TGA data of polyorganosiloxane polymers and composites. Temperature for mass loss

Sample

15%

50%

PDMS pBSil PBSil–OSEP(5%) PBSil–MB(5%) PBSil–OSEP(5%)–MB(5%)

398 415 424 431 421

542 488 498 476 477

Temperature for maximum rate of mass loss (1C) 352 – – – –

Residue at 550 1C (%)

509 481 489 472 475

48 38 36 33 30

step volatile products, mostly cyclic compounds, evolve and fragments of lower molecular weight are formed. In the second step further depolymerization and breakdown of the PDMS chain takes place. In the oxidative degradation of pBSil with respect to that of PDMS, the first degradation step completely disappears. The residue at 550 1C is smaller for pBSil, probably because boron can form volatile compounds, which are easily lost.13 The characteristic mass-loss data obtained in air atmosphere are presented in Table 4.2. The boron coupled to the polysiloxane chain of pBSil forms a network through borosiloxane linkages. The temporary B


O interchain dative bonds in polyborosiloxane also establish the linkages of a network structure. The incorporated boron favours reactions such as Si–CH3 dehydrogenation to Si–CH2 and the splitting off the methyl groups from PDMS through conversion into radicals and giving additional binding sites, as shown in Equations 4.1–4.3.14  Si  CH3 ! Sid þ dCH3

ð4:1Þ

 Si  CH3 þ dCH3 ! Si  CH2 d þ CH4

ð4:2Þ

 Sid þ dCH2  Si ! Si  CH2  Si 

ð4:3Þ

A significant enhancement in thermal stability and the disappearance of the first degradation step are observed in TG curves at low temperature (see Figure 4.1). This can be ascribed to the cross-linking structure of pBSil, which hinders the splitting of the cyclic oligomers and the formation of disadvantageous volatile products. This effect is more pronounced in the presence of MB and OSEP.

4.3.2 Combustion Characteristics of Polypropylene-Based Composites The pBSil–OSEP(5%)–MB(5%) sample was selected to prepare PP-based composites. The samples made of PP and polyborosiloxane composites are summarised in Table 4.3, where the values are given in weight percent.

54

Table 4.3

Chapter 4

The compositions of PP–(pBSil–OSEP–MB) in wt%. PBSil–OSEP–MB

Sample code NKMB060 NKMB061 NKMB062 NKMB063 NKMB073

PP 0 75 60 45 45

PBSil–OSEP–MB a

100 25 40 55 55*

pBSil

OSEP

MB

90 22.5 36 49.5 49.5

5 1.25 2 2.75 2.75

5 1.25 2 2.75 2.75

a

0.5% peroxide content.

Figure 4.2

Rate of heat release vs time curves of PP–(pBSil–OSEP–MB) composites as the percentage of pBSil–OSEP–MB was varied from 0 to 55 wt% (see Table 4.3); irradiation heat flux: 50 kW m2.

The HRR curves of the composites are shown in Figure 4.2. The amount of pBSil–OSEP–MB has a significant influence on the combustion properties. Peak HRR decreases with increasing percentage of pBSil–OSEP–MB in the composite and, furthermore, reduction in THR and effective heat of combustion (EHC) can also be observed. However, the TTI time range decreases slightly. All the quantitative results are given in Table 4.4. The lower HRR (kW
m2) and THR (MJ
m2), indicate a smaller fire hazard, but with a slightly shortened TTI (s). The curves of the PP–(pBSil–OSEP–MB composites lie between the two references, pure PP and pBSil–OSEP–MB. The shape of the curves changes

55

Use of Organosilicone Composites as Flame Retardant Additives

Table 4.4 Sample PP NKMB060 NKMB061 NKMB062 NKMB063 NKMB073

Figure 4.3

Combustion properties of PP/pBSil-OSEP-MB composites. TTI (s)

peak HRR (kW
m2)

34 22 24 25 20 20

898 132 586 441 332 230

Peak time (s) 149 62 208 227 263 256

THR (MJ
m2) 115.3 23.3 84.1 76.3 67.1 43.8

Residue (%) 0 19.8 2.3 3.7 6.1 7.9

Images of the combustion residue of (a) PP and (b) PP–(pBSil–OSEP– MB) after treatment under a cone heater (heat flux: 50 kW m2). Image (c) shows the non-carbonaceous, ceramified intumescent structure.

gradually with increasing amounts of pBSil–OSEP–MB. For pristine PP, a shoulder occurs before the maximum HRR value, while as the pBSil content of the composites is increased this shoulder gradually broadens. The presence of pBSil–OSEP–MB causes a steady-state heat release until the peak at the end of the combustion, which also decreases. Applying peroxide as radical initiator (sample NKMB073), the thermal stability and the combustion properties improved further. As shown in Table 4.4, the HRR peak reduced from 898 to 230 kW m2. The time at which this occurred increased from 146 to 256 seconds, and a remarkable decrease occurred in the THR value. Interestingly, at the beginning of combustion the composite NKMB073 exhibits a HRR as low as that of the pure pBSil–OSEP–MB reference. The peroxide probably acts by building linkages between polymer chains both within the silicone phase and at the interfaces of the two polymers, but validation of this assumption requires further experiments. The advantageous effect of the pBSil systems in PP can be explained by the formation of a non-carbonaceous, white, ceramified, intumescent structure on the surface during the combustion, which acts as a protective ceramic layer. At the end of combustion the PP leaves no residue, whereas the PP–(pBSil–OSEP– MB) samples leave a white solid coherent char-like residue. The thickness of the residue is about 3–5 centimetres, as shown in Figure 4.3.

56

Chapter 4

4.3.3 Multilayer Structure PP–(pBSil–OSEP–MB) The efficiency of the pBSil–OSEP–MB composite as a multilayer structure was also investigated, using it as a coating on the PP surface. In previous studies on boron-containing PDMS the authors found that the boron content helps to form a protective, coherent, heat-resistant ceramic layer on the polymer during combustion.4,15 These results initiated the using of polyborosiloxane composite as a coating on the PP surface. During processing we found that the pBSil–OSEP–MB composite readily forms a continuous layer on the surface of PP. The 500 mm layer formed on the PP core of 4 mm corresponds to less than the half of the lowest pBSil–OSEP– MB concentration (NKMB061) in Table 4.3. Such multilayer structures resulted in improved combustion characteristics (see Figure 4.4) – more than a 30% reduction in peak HRR and a considerably delayed peak (from 217 to 454 seconds) compared to PP. Furthermore, in this case the TTI did not decrease, but increased by 60% of the original value. The wettability study of this surface revealed that the coated layer is less adhesive than that of pristine PP, as shown in Figure 4.5. A digital camera is mounted on the contact angle meter. The contact angle of water drop on the PP surface is 901 while the coated surface has a contact angle value of 1151. The results suggest that multilayer structures may provide protection more efficiently, especially against the early degradation of the polymer, than mixtures of the same composition. Such surface layers may perform multifunctional

Figure 4.4

HRR results of the reference polymer and the multilayer structure; irradiation heat flux: 50 kW m2.

Use of Organosilicone Composites as Flame Retardant Additives

Figure 4.5

57

Photo of water drop on (a) PP coated with pBSil–OSEP–MB material and (b) uncoated PP as reference.

roles, acting both as flame retardant and as an adhesion–modifying (dust and/or graffiti releasing, self-cleaning) layer.

4.4 Conclusion The thermo-oxidative degradation of PDMS is affected by the inclusion of boron atoms. The pBSil formed by a polycondensation reaction exhibits higher stability in the initial stage of decomposition as the boron-containing units hinder the splitting of the cyclic oligomers. At higher temperatures, however, the mass loss of pBSil is larger than that of PDMS. To improve the fire retardancy of PP, pBSil was combined with MB and a fibrous clay (sepiolite). Increasing the concentration of pBSil–OSEP–MB in PP decreases the peak HRR gradually, as well as the THR value. These results suggest that the formation of a non-carbonaceous, white, ceramified, intumescent layer on the surface of PP may act similarly to the conventional carbonaceous foam, but its heat stability is greater. Further improvement was achieved by the introduction of a radical initiator into the system. This probably acts by building linkages between the polymer chains. Also, pBSil–OSEP–MB readily forms a continuous layer on the surface of PP. Applied as a surface layer it acts as a fire-retardant protecting coating and also enables the control of adhesion at the surface. This double effect of the developed coating is advantageous in construction, automotive and some other relevant industries.

Acknowledgements Supports from EU 6 Multihybrids (IP 026685-2), Nanofire (NMP3-CT 2004505637) projects, Hungarian Research Fund OTKA T049121, Fund of

58

Chapter 4

European Union and Hungarian State GVOP/3.1.1.-2004-0531/3.0, Public Benefit Association of Sciences and Sport of the Budapest University of Technology and Economics are acknowledged.

References 1. P.J. Davies and A.J. Fletcher, J. Mech. Eng. Sci., 1995, 209(6), 408. 2. A. Juhasz, P. Tasnadi and L. Fabry, Phys. Educ., 1984, 19, 302. 3. I. Ravadits, A. To´th, G. Marosi, J. Papp and S. Szabo´, Polym. Degrad. Stabil., 2001, 74, 419. + and A. Sze´p, 4. G. Marosi, A. Ma´rton, P. Anna, G. Bertalan, B. Marosfoi Polym. Degrad. Stabil, 2002, 77, 259. 5. A. Toldy, P. Anna, I. Csontos, A. Szabo´ and G. Marosi, Polym. Degrad. Stabil, 2007, 92(12), 2223. 6. http://www.borax.com. 7. S. Bourbigot, M. Le Bras and S. Duquesne, in Fire Retardancy of polymers, eds. M. LeBras, S. Bourbigot, S. Duquesne, C. Jama, C. Wilkie, The Royal Society of Chemistry, Cambridge, 2005, p. 327. 8. F. Samyn, S. Bourbigot, S. Duquesne and R. Delobel, Thermochim. Acta, 2007, 456, 134. 9. A. Dechirico, G. Audisio, F. Provasoli, M. Armanini and R. Franzese, Macromol. Symp., 1993, 74, 343. 10. A.B. Morgan, J.L. Jurs and J.M. Tour, J. Appl. Polym. Sci., 2000, 76(8), 1257. 11. G. Camino, S.M. Lomakin and M. Lageard, Polymer, 2002, 43, 2011. 12. W. Zhou, H. Yang, X. Guo and J. Lu, Polym. Degrad. Stabil., 2006, 91(7), 1471. 13. G.D. Soraru, F. Babonneau, S. Maurina and J. Vicens, J. Non-Cryst. Solids, 1998, 224, 173. 14. R. Pen˜a Alonso, F. Rubio, J. Rubio and J.L. Oteo, J. Anal. Appl. Pyrol., 2004, 71, 827. 15. G. Marosi, P. Anna, A. Ma´rton, G. Bertalan, A. Bo´ta, A. To´th, M. Mohai and I. Ra´cz, Polym. Advan. Technol., 2002, 13, 1103.

CHAPTER 5

Organomodified Ultrafine Kaolin for Mechanical Reinforcement and Improved Flame Retardancy of Recycled Polyethylene Terephthalate B. SWOBODA,a E. LEROY,a, c J.-M. LOPEZ CUESTA,a C. ARTIGO,b C. PETTERb AND C.H. SAMPAIOb a

Ecole des Mines d’Ale`s, Centre des Mate´riaux de Grande Diffusion (CMGD), 6 Av. de Clavie`res, 30319 Ale`s Cedex, France; b Laborato´rio de Processamento Mineral (LAPROM), Centro de Tecnologia – UFRGS, Av. Bento Gonc¸alves, 9500, Caixa Postal 15021-91501-970, Porto Alegre, Brasil; c Current address: Laboratoire de Ge´nie des Proce´de´s-EnvironnementAgroalimentaire (GEPEA), 37, Bd de l’Universite´, 44606 Saint-Nazaire, BP 420 Cedex, France

5.1 Introduction The development of polymer composites that contain ultrafine, delaminated or exfoliated phyllosilicates is increasing everyday. Minerals of high aspect ratio provide large interfacial areas between the mineral particles and polymer chains, which result in significant improvements in the tensile and flexural mechanical properties of the polymer composite. Barrier properties can also be

Fire Retardancy of Polymers: New Strategies and Mechanisms Edited by T Richard Hull and Baljinder K Kandola r Royal Society of Chemistry 2009 Published by the Royal Society of Chemistry, www.rsc.org

59

60

Chapter 5

conferred by the use of phyllosilicates in thermoplastics, which significantly influences mass transfer (limit gases and solvent permeation) and the reaction to fire because a carbonaceous and mineral protective layer caused by polymer ablation forms during combustion.1,2 Organomodified montmorillonites are the main category of phyllosilicates used in polymer composites to improve the above-mentioned properties, since they are able to intercalate with polymer chains or exfoliate inside the polymer matrix.1,2 Phyllosilicates other than organomodified montmorillonites, such as kaolinites, sepiolites attapulgites etc., are advantageous since it is possible to carry out functionalisation, in particular by reaction with the hydroxyl groups. This is not easy with montmorillonite because the hydroxyl groups are not available. Kaolinite is a 1 : 1 phyllosilicate with the general chemical formula Al2Si2O5(OH)4.3 Contrary to montmorillonite, kaolinite does not have interlayer exchangeable cations. Nevertheless, the external surface of kaolinite particles is covered with silanol groups, which are able to react with organic molecules. Moreover, depending on the origin of the mineral deposit and on the mineral processing conditions, ultrafine kaolinite can be obtained with submicronic average particle size and high aspect-ratio particle morphology, which make it suitable for the nanoscale reinforcement of a polymer matrix and improvement in barrier properties. We used unmodified kaolinite, kaolinite modified with triphenylphosphite (TPP) and a physical mixture of kaolinite and TPP to study the effect of organic modification on dispersibility in a recycled polyethylene terephthalate (PET) and, consequently, its effect on the mechanical properties and reaction to fire of the polymer composite. Recycled PET was chosen because of its intrinsically low mechanical performance and thermal stability. It is also a technological challenge to be able to upgrade the relevant properties of recycled PET compounds for use in building sector and electrical equipment. TPP was chosen for the organic modification of kaolinite for two reasons: first, its phosphorus content means it is likely to act as a flame retardant for polyesters through enhanced charring.4 This is particularly important for PET since this polyester has a particularly poor charring behaviour. Moreover, TPP is a known ‘‘chain extender’’ for PET. In the melt, the presence of TPP leads to a bridging reaction between PET chain ends [Equation (5.1), in which R1 and R2 are PET chains and Ph is a phenolic ring], which has been studied in detail by various authors,5–7 and leads to a significant increase in the average molecular mass and mechanical properties of the polymer. R1
OH þ R2
COOH þ P
ðOPhÞ3 ! R1
OCO
R2 þ HO
P
ðOPhÞ2 þ PhOH

ð5:1Þ

In the following sections, various PET compounds are characterized in order to compare the efficiency of TPP-modified kaolinite particles (for improving

Organomodified Ultrafine Kaolin for Mechanical Reinforcement

61

mechanical and reaction to fire properties) to that of unmodified kaolinite particles, ungrafted TPP and their mixtures.

5.2 Experimental 5.2.1 Materials Three different kaolin samples (named C, P and D) were obtained from different mineral deposits located in Amazonia using a pilot plant unit of the Vale do Rio Doce Company (Brazil) and were submitted to different mineral-processing operations, including size selection. TPP was purchased from Acros Organics (CAS: 101-02-0). The recycled PET (post-consumer bottles flakes) was supplied by VALORPLAST, France, (ZPET ¼ 0.76 dl g
1, in 2-chlorophenol at 25 1C, average molecular weight ¼ 26 300 g mol
1, calculated using the Mark Houwink equation (K ¼ 3.8  104; a ¼ 1.3).

5.2.2 Processing 5.2.2.1

Organic Modification of Kaolin

Kaolinite D powder was placed in a reactor at 80 1C and TPP (liquid) was progressively introduced with vigorous stirring until a kaolinite : TPP ratio of 70 : 30 (w/w) was achieved. The reactor was then sealed to avoid the evaporation of TPP. Samples were taken from the reactor after 24 hours and 48 hours for analysis. The wet powder was filtered after 72 hours of stirring, washed with tetrahydrofuran (THF) and dried at 80 1C for 12 hours. It was then used for melt processing with PET.

5.2.2.2

Polymer Compounding

The high sensitivity of PET to moisture meant we used the following protocol before melt mixing: PET flakes were dried in vacuum at 120 1C for 16 hours. These conditions were chosen on the basis of a detailed study of PET drying kinetics and sensitivity to moisture during melt processing.8 Four polymer compounds were prepared by extrusion followed by injection moulding, all containing a mineral : additive loading of 5weight percent (wt%) in mass:  PET–unmodified Kaolin D (95 : 5 w/w), called PET/K herein,  PET–TPP (95 : 5 w/w), called PET/TPP herein,  PET–unmodified Kaolin D–TPP (95 : 3.5 : 1.5 w/w), called PET/K/TPP herein,  PET–organomodified Kaolin D (95 : 5 w/w), called PET/MK herein.

62

50

100

T(°C)

25 25

200

275

150

50

50

K2

270

25

100

50

K1

260

25 50

80

Chapter 5

length (mm)

12.5

12.5

1200 mm

Legend

Thread : 33 mm

Kneading disc blocs : 90°

Thread : 25 mm

Kneading disc blocs : 45°C

Thread : 16 mm

Reverse thread : 25 mm

Conveying screw elements

Figure 5.1

Kneading screw elements

Screw configuration used for extrusion.

Pure recycled PET reference samples were prepared using the same processing conditions: blending was performed in a twin-screw co-rotating extruder Clextral BC 21 (L ¼ 1200 mm, L/D ¼ 48). Figure 5.1 shows the screw configuration designed to promote kaolin dispersion. Extrusion conditions were same for all compositions: feed rate (Q) ¼ 6 kg h
1 and screw speed (N) ¼ 350 revolutions per minute (rpm). The residence time in the extruder under these conditions was about 75 seconds. Granulation was performed after extrusion. Before injection moulding, the granules were dried in an oven at 105 1C overnight at reduced pressure [approximately 1 kPa (10 mbar)]. Injection moulding was carried out at 250–265 1C using a Sandretto Serie Otto AT 95 machine with a mould temperature of 40 1C. The specimens produced were both dog bones (according to ISO 527-2 type 1A specifications for mechanical testing) and 100  100  4 mm sheets (used for fire testing).

5.2.3 Characterization Techniques 5.2.3.1

Characterization of Unmodified Kaolins

Particle-size measurements were performed on a Cilas Laser apparatus based on light diffraction, and the specific surface area was determined using the Brunauer– Emmett–Teller (BET) method by nitrogen absorption at 77 Kelvin (
183 1C). The complete chemical composition of the samples was determined by X-ray fluorescence. The mineralogical study was carried out using a Siemens D-50 X-ray diffractometer on non-orientated samples (powder method). Finally, the particle size, and morphological and textural features of the kaolins were observed using a Quanta 200 FEG environmental scanning electron microscope (FEI Company).

Organomodified Ultrafine Kaolin for Mechanical Reinforcement

5.2.3.2

63

Characterization of Organomodified Kaolin

The percentage of TPP that remained on the kaolin powder after washing and drying was determined by calcination at 700 1C in air. In parallel, the thermo-oxidative degradation of modified kaolin was studied by thermogravimetric analysis (TGA; 5 1C min
1, 25–700 1C, samples of typically 25 mg, using a Netzsch STA 409 device) and compared to that of pure TPP and unmodified kaolin D. As for unmodified kaolin, X-ray diffraction was performed on non-orientated samples (powder method). The electrophoretical mobility (zeta potential) and the average particle sizes of pristine and modified kaolin D were measured in water and THF, using a Malvern Nanosizer NANOZS, to characterize the organophilic behaviour of modified kaolin. Finally, the grafting reaction of TPP onto a kaolin surface was investigated by Fourier transform infrared (FTIR) spectroscopy using a Bruker IFS66 spectrometer ATR mode (32 scans, resolution of 4 cm
1).

5.2.3.3

Characterization of Polymer Compounds

To study the interactions between the different components during melt mixing, blends were made using a Haake Rheomix internal mixer (in parallel to the extrusion processing) and the torque was measured as function of time (T ¼ 270 1C and rotor speed ¼ 60 rpm). PET was introduced in the mixer prior to the introduction of the other components. The morphology of polymer compounds obtained from the extrusion–injection process was observed using a Quanta 200 FEG environmental scanning electron microscope (FEI). The dispersion of kaolin particles in the PET matrix was also studied indirectly by performing dynamic rheological tests at a low shear rate at 260 1C using an ARES Rheometrics Scientific apparatus in plate–plate geometry (+ 25 mm, 1 mm gap). Flexural modulus was measured from three-point bending tests performed on a Zwick TH010 universal press according to ISO 178 Standard. TGA (Setaram) was performed under air flow (75 ml min
1) using samples of typically 30 mg in platinum pans, submitted to a temperature ramp from 25 1C to 700 1C at a heating rate of 5 1C min
1. The limiting oxygen index (LOI) was measured using a Stanton Redcroft instrument on barrels (80  10  4 mm3), according to ISO 4589 specifications. Cone calorimeter tests (ISO 5660) were performed on 100  100  4  mm3 samples placed horizontally, with an irradiance of 50 kW m
2. Only timeto-ignition (TTI), peak heat-release rate (pHRR) and the fire propagation index (FPI ¼ pHRR/TTI) values are discussed below. Results correspond to mean values obtained from three experiments for each formulation, for which a typical variation of 10% was observed. The charring behaviour of the different formulations was evaluated by placing 1 g samples at 700 1C for 4 minutes in an oven under air atmosphere. After this combustion, the mass fraction of charred residue was measured.

64

Chapter 5

Table 5.1

Chemical composition of kaolins.

Sample SiO2 Al2O3 TiO2 Fe2O3 C P D

45.4 45.1 44.9

40.0 39.7 39.1

0.63 0.89 1.6

0.55 0.54 1.1

MnO

MgO CaO Na2O K2O P2O5

o0.01 o0.1 0.01 o0.01 0.11 0.09 o0.01 0.15 0.01

0.27 0.27 0.30

0.02 0.02 0.02

0.05 0.09 0.05

PF 14.02 14.05 14.45

5.3 Results and Discussion 5.3.1 Properties of Unmodified Kaolins The chemical composition of each sample as a percentage of oxide present is given in Table 5.1. Kaolin D presents the highest level of contamination, with Fe and Ti concentrations of 1.1 and 1.6%, respectively. According to X-ray diffraction experiments, titanium dioxide in the form of anatase was the only mineralogical type associated with kaolinite, particularly for kaolin D. Figure 5.2 shows representative scanning electron microscope (SEM) micrographs of the three kaolin samples. Table 5.2 gives their particle-size distribution – the percentages correspond to volumetric concentrations of particles in determined sized ranges:  Kaolin C is characterized by well-crystallized kaolin platelets, with rectilinear edges; it presents two well-defined populations. The average size of particles is in the range 0.4–2.3 mm.  Kaolin P is composed of kaolinite platelets with an average size in the range 0.4–0.7 mm. There is also a significant population of ultrafine particles with an average size of approximately 250 nm.  Kaolin D is an ultrafine kaolin composed entirely of isolated kaolinite platelets of euhedric form with average size in the range 200–400 nm. The results obtained from BET measurements of specific surface area reveal a direct correlation with the average size of the kaolin particles. Kaolin D presents a significantly higher specific surface area (21.8 m2 g
1) than kaolins C and P, with specific surface areas of 9.9 and 12.1 m2 g
1, respectively. From these results, kaolin D was selected for modification by TPP and compounding with recycled PET.

5.3.2 Grafting of TPP onto Kaolin Surface Modified kaolin D samples obtained after treating with TPP for 72 hours, washing with THF and drying (see Section 5.2.2.1), contained 30 wt%. of organic modifier (as measured by calcination). The TGA of modified kaolin D is presented in Figure 5.3, along with those of pure TPP and unmodified kaolin D. In the temperature range studied, the unmodified kaolin only shows a first stage of mass loss below 100 1C, which corresponds to physically adsorbed

Organomodified Ultrafine Kaolin for Mechanical Reinforcement

Figure 5.2

SEM micrographs of kaolins (top to bottom: C, P and D), 50 000.

65

66

Chapter 5

Granulometric distribution of kaolins.

Table 5.2 Sample

%, o10 mm

%, o5 mm

%, o2 mm

%, o1 mm

%, o0.5 mm

%, o0.1 mm

D50 (mm)

C P D

98.52 100.00 100.00

88.64 100.00 100.00

59.66 90.66 96.03

48.32 73.74 90.35

33.30 49.19 70.25

4.58 7.26 16.29

1.13 0.51 0.32

100

Kaolin

90 80

Modified Kaolin

Mass loss (%)

70 60 50 40 30 20

TPP

10 0 0

100

200

300

400

500

600

Temperature (°C)

Figure 5.3

Thermogravimetric analysis of kaolin, TPP and modified kaolin.

water. Thereafter, the mineral is stable up to around 400 1C, which corresponds to the beginning of the dehydroxylation followed at higher temperatures by transformation into metakaolinite.3 Pure TPP and modified kaolin D both exhibit a mass loss starting around 150 1C. While pure TPP nearly reaches complete degradation, the mass loss of modified kaolin reaches a plateau at 70% of residual mass, which is in agreement with the estimate of organic modifier content by calcination. The lack of significant differences between the degradation temperatures of pure TPP and modified kaolin suggest that TPP was not intercalated between kaolin platelets. This proposition is supported by X-ray diffraction measurements, which showed that the spacing between kaolin layers remained constant at 7.3 A˚, which is the characteristic feature of pure kaolinite.9 Nevertheless, electrophoretic and granulometric measurements in water and THF (Table 5.3) clearly show that the surface properties of kaolin have been affected by the treatment, such that the unmodified kaolin shows smaller particulate size and higher Zeta potential in water than in THF. This can be attributed to the

67

Organomodified Ultrafine Kaolin for Mechanical Reinforcement

Table 5.3

Zeta potential and average particle size of kaolin and modified kaolin in water and THF solvents. Water

Solvent Average size (mm) Zeta potential (mV)

THF

Kaolin

Modified kaolin

Kaolin

Modified kaolin

0.32 40.9

2.79 0.2

0.98 18.2

0.85 48.1

hydrophilic nature of untreated kaolin particles’ surface. On the contrary, the zeta potential of modified kaolin in water is very small, and both this and a significantly increased average particle size indicate the aggregation of individual particles and thus a hydrophobic behaviour of particle surfaces. The zeta potential of modified kaolinite increased and the average particle size decreased in THF, indicating an organophilic behaviour. The FTIR analysis of kaolin powder, before modification and after 24, 48 and 72 hours of treatment with TPP (followed by THF washing and drying) shows two important changes in the absorption spectra of the modified kaolin (Figure 5.4):  a peak appears at 1193 cm
1 and can be attributed to the formation of Si–O–P chemical bonds.10  In parallel, there is an attenuation of the peaks between 3600 and 3750 cm
1 which corresponds to surface hydroxyl groups of kaolin.11 Given that the reactivity of kaolin surface hydroxyl groups is similar to that of a carboxylic acid function,12 the following reaction mechanism can be assumed: kaolin
Si
OH þ P
ðOPhÞ3 ! kaolin
Si
O
P
ðOPhÞ2 þ PhOH ð5:2Þ

5.3.3 Morphological, Rheological and Mechanical Properties of Polymer Compounds Figure 5.5 shows the torque evolution during the melt mixing in the internal mixer. The composition that contained pristine kaolin (PET–K) exhibits low torque values, comparable to those obtained for pure PET. These low values can be ascribed to the promotion of PET chain breaking by the surface hydroxyl groups present at the kaolin surface. Conversely, the PET–TPP composition has a high viscosity because of the chain-extension reaction. When both additives are introduced (sample PET–K–TPP), the influence of TPP dominates, but with much slower kinetics. Finally, for modified kaolin (PET– MK), two steps can be seen. First, a chain-extension effect seems to occur, which shows that a part of TPP fixed onto kaolin was physically adsorbed and is able to react with PET. Then, the torque increases more slowly. That the torque does not decrease suggests that the presence of TPP on the kaolin

68

Chapter 5

Transmittance (normalized)

after 72h treatment

after 48h treatment

after 24h treatment

untreated Kaolin D

700

800

900

1000 1100 Wave number (cm-1)

1200

1300

Transmittance (normalized)

after 72h treatment

3400

Figure 5.4

after 48h treatment

after 24h treatment

untreated Kaolin D

3500

3600

3700 3800 Wave number (cm-1)

3900

4000

FTIR analysis of kaolin modification by TPP.

surface (replacing surface hydroxyl groups) could limit the chain-breaking effect of kaolin. Figure 5.6 shows SEM micrographs obtained for PET–K and PET–MK compounds. On the micrographs on the left, in both cases the kaolin particles are homogeneously dispersed in the PET matrix. However, closer examination at higher magnification (micrographs on the right in Figure 5.6) reveals that in

69

Organomodified Ultrafine Kaolin for Mechanical Reinforcement 110 100 90

Torque (N.m)

80 70

PET

60

PET/K

50

PET/TPP

40

PET/K/TPP

30

PET/MK

20 10 0 0

Figure 5.5

2

4

6 mixing time (min)

8

10

12

Evolution of torque during melt mixing of polymer compounds (vertical scale corresponds to the pure PET sample, the other curves have been shifted for the sake of clarity).

the case of pristine kaolin, the aggregates are significantly bigger than those for modified kaolin. This indicates that the organophilic modification promotes the dispersion of individual kaolin platelets in the polymer. The bending modulus values are presented in Table 5.4. The best results are shown by the PET–MK sample, with a 60% increase in bending modulus compared to that of the recycled PET sample. The improved kaolin-reinforcement effect in PET–MK is also significantly better than that of pristine kaolin with the latter showing a 40% increase in bending modulus compared to pure PET. The improved mechanical performance of PET–MK with a lower mineral loading of 3.5 wt% suggests better dispersion and better compatibility of kaolin platelets with the PET matrix. It is also possible that reactions between grafted TPP and PET could have occurred in the melt, leading to chain extension and/ or coupling reactions between kaolin particles and PET chains. However, this is beyond the scope of this investigation and will be the subject of future work. PET–TPP and PET–K–TPP compounds also show significant improvements in mechanical properties, which could be due to PET chain extension and/or nanoscale reinforcement of the PET matrix by kaolin particles.

5.3.4 Thermal Stability and Reaction to Fire of Polymeric Compounds The mass-loss curves obtained in TGA experiments (Figure 5.7) show that all the samples exhibit a two-step degradation: at lower temperatures, the thermooxidation of the polymer matrix leads to both low molecular weight compounds

70

Figure 5.6

Table 5.4

Chapter 5

SEM micrographs of compounds (a) PET–K and (b) PET–MK.

Mechanical and fire performance of different polymeric compounds.

Sample

PET

PET–TPP

PET–K

PET–K–TPP

PET–MK

Bending modulus (MPa) LOI (0.5%) Cone Calorimeter TTI (s) pHRR (kW m
2) FPI (kW m
2 s
1) THR (MJ m
2) Char residue after 4 min at 700 1C in air (wt %)

2017 (272) 22.5

2866 (26)

2999 (136)

27.2

2862 (41) 26.4

26.6

3313 (86) 29.7

71 710

93 502

91 517

77 352

97 312

10.0

5.4

5.7

4.6

3.2

99 0.0

49 12.0

44 3.5

60 5.5

69 9.5

71

Organomodified Ultrafine Kaolin for Mechanical Reinforcement 100

Mass loss (%)

80

PET PET/TPP PET/K PET/K/ TPP PET/MK

60

40

20

0 330

350

370

390

410

430

450

470

490

510

530

550

570

590

T (°C)

Figure 5.7

Thermogravimetric analysis of the different polymer compounds.

and char formation. The first mass loss observed corresponds to the evaporation of these low molecular weight compounds, while the second mass loss corresponds to the degradation of the char previously formed at lower temperatures. The PET–MK compound clearly exhibits superior thermal stability, with the two mass-loss stages shifted towards higher temperatures compared to pure PET. This can be correlated with a significant increase of both TTI (+30%) in cone calorimeter tests and char residue in oven tests (Table 5.4). However, the presence of TPP has accelerated the initial mass loss compared to that of pure PET. However, enhanced char formation at higher temperatures (4500 1C) occurs in the presence of TPP. This observation can be related to the limited thermal stability of TPP, which starts to decompose at around 200 1C (see Figure 5.3) and to its flame retardant mode of action, which promotes char formation. It can be proposed that the decomposition of TPP leads to species that accelerate chain-breaking reactions, which results in the lower thermal stability of PET at lower temperatures, followed by cross-linking reactions that promote char formation at higher temperatures. The presence of pristine kaolin also accelerates the initial mass loss, probably through the promotion of chainbreaking reactions by surface hydroxyl groups on kaolin platelets. The dehydroxylation of kaolin can promote PET hydrolysis. When TPP-grafted kaolin particles are used (PET–MK), this proposed effect of surface hydroxyl groups on PET degradation is obviously limited. Nevertheless, the accelerations of the first mass loss for PET–TPP, PET–K and PET–K–TPP are not correlated with a decrease in TTI since all these compounds present higher values than pure PET (Table 5.4). The lack of a direct correlation between the thermal stability of the PET matrix in TGA experiments and the TTI in cone calorimeter tests is not surprising for such

72

Chapter 5 800 700

PET PET/K

HRR (kW/m2)

600

PET/TPP PET/K/TPP

500

PET/MK 400 300 200 100 0 0

100

200

300

400

500

600

time (s)

Figure 5.8

Typical heat release rate vs. time curves obtained in cone calorimeter experiments.

systems since additional phenomena able to retard ignition have to be taken into account. In particular, barrier effects to fuel transfer from the condensed to gaseous phases caused by char formation and/or kaolin platelets. Regarding other aspects of the reaction to fire (Table 5.4 and Figure 5.8) a significant difference can be observed between PET–MK and other samples – the former presents better performance for most of the tests. Compared to pure PET, LOI increases by 30% and pHRR decreases by 50%. However, the total time to burn has increased, hence total heat release (THR) is not much affected and it is not the best for PET–MK. For all that, it has to be taken into account that for PET–TPP and PET–K the loadings of TPP (and the associated char yield) and mineral, respectively, are higher. In addition, PET–MK offers the best combination between THR and FPI reductions. It is likely that kaolin platelets act as a reinforcement for the char formed, as observed in previous studies on recycled PET.13,14: when PET is flame retarded by systems based on red phosphorus and alumina or iron oxide, it has been shown that the introduction of talc particles or glass fibres could reinforce and stabilize the char formed, leading to a better reaction to fire. When TPP-grafted kaolin particles are used (PET–MK), the promotion of charring is strongly enhanced compared to that of PET–K–TPP (Table 5.4), even though both compounds contain the same amounts of mineral and organic additives. It can be assumed that the presence of well-dispersed modified kaolin platelets in the PET matrix leads to barrier properties that limit the evaporation of TPP and its degradation products above 200 1C. Consequently, less TPP (and degradation products) escapes to the gas phase and more remains in the condensed phase to act as a char promoter.

Organomodified Ultrafine Kaolin for Mechanical Reinforcement

73

5.4 Conclusion That ultrafine kaolins present BET surfaces up to 20 m2 g
1 as fillers for polymers, as highlighted here, is of interest. The possibility to easily modify the hydrophilic behaviour of kaolin particles to organophilic behaviour has been shown. No intercalation of polymer between the kaolin platelets was involved, contrary to the modification of bentonite or montmorillonite clays, but grafting reactions with surface hydroxyl groups of kaolin platelets were evident. In the system studied, this grafting reaction was carried out in bulk, by directly mixing the kaolin with the organic modifier. The organic modification of kaolin platelets by TTP (an additive able to act both as a char promoter and a chain extender in PET) was shown to improve their dispersion in PET and to limit the chain-breaking effects of surface hydroxyl groups (present in pristine kaolin surfaces). Moreover, incorporation of modified kaolin, at a loading level of only 5 wt%, in recycled PET significantly increased the mechanical and reaction-to-fire properties of the recycled PET.

5.5 Acknowledgements The authors thank the CAPES COFECUB for supporting this Franco-Brazilian research programme.

References 1. 2. 3. 4. 5. 6. 7. 8. 9. 10. 11. 12. 13. 14.

J.W. Gilman, Appl. Clay Sci., 1999, 15, 31. M. Alexandre and P. Dubois, Mater. Sci. Eng.: R: Reports, 2000, 28, 1. R.E. Grim, Applied Clay Mineralogy, Mc Graw Hill, New York, 1962. S.V. Levchik, ED Weil. Polym. Int., 2005, 54, 11. S.M. Aharoni, C.E. Forbes, W.B. Hammond, D.M. Hindenlang, F. Mares, K. O’Brien and R.D. Sedgwick, J. Polym. Sci. Part A, 1986, 24, 1281. B. Jacques, J. Devaux, R. Legras and E. Nield, Polymer., 1997, 38, 5367. B. Jacques, J. Devaux, R. Legras and E. Nield, Polymer., 1996, 37, 5367. B. Swoboda, E. Leroy, J.-M. Lopez-Cuesta, Special Issue MoDeSt’06, Polym. Deg. Stab., 2007. F. Franco and M.D. Ruiz Cruz, Clay Minerals, 2004, 39, 193. B.A. Morrow and S.J. Lang, J. Phys. Chem., 1994, 98, 13319. R.L. Frost, J. Colloid Int. Sci., 2002, 1246, 164. Mineral Fillers for Rubber : kaolin clay, RT Vanderbilt Company, (2007). F. Laoutid, L. Ferry, J.M. Lopez Cuesta and A. Crespy, Fire Mater., 2006, 30, 343. F. Laoutid, L. Ferry, J.M. Lopez Cuesta and A. Crespy, Polym. Deg. Stab., 2003, 82, 357.

CHAPTER 6

Complex Micro-analysis Assisted Design of FireRetardant Nanocomposites – Contribution to the Nanomechanism + P. ANNA AND GY. MAROSI A. SZABO´, B. B. MAROSFOI, Department of Organic Chemistry and Technology, Budapest University of Technology and Economics, 1521 Budapest, Hungary

6.1 Introduction Advancement of the methods used to produce fire retardancy is hindered by the shortcomings of the current characterization methods and insufficient understanding of the fire retardancy mechanisms. These two factors are connected with each other: better characterization would promote better understanding and mechanistic studies initiate advancements in analysis. A variety of methods is available for characterization, but none of them provide all the data required to model the combustion process adequately. Development is needed, especially of the smalland microscale analytical methods, to allow rapid screening of small experimental samples and possibly to enable high-throughput and/or in-line characterization. D. Price was the pioneer who realized, earlier than anyone else the importance of very small-scale modelling in fire testing, which became essential when the

Fire Retardancy of Polymers: New Strategies and Mechanisms Edited by T Richard Hull and Baljinder K Kandola r Royal Society of Chemistry 2009 Published by the Royal Society of Chemistry, www.rsc.org

74

Complex Micro-analysis Assisted Design of Fire-Retardant Nanocomposites

75

1,2

expensive nanofillers appeared. He developed a laser pyrolysis time-of-flight mass spectroscopy (LP-ToFMS) system to understand the fundamentals of fire retardancy through analysis of the gases that evolved during a mini-fire induced by a laser with well-controlled energy.2,3 This elegant technique inspired the authors, being interested in condensed-phase fire retardant activity, to extend the method to the analysis of the condensed phase using a LP micro-Raman system.4 The Raman method is favoured at the set-up of such systems because of its applicability for in-depth measurements, its good lateral and in-depth resolution, its unique sensitivity to double bonds and physical (e.g. polymorphic) changes, and its insensitivity to water. If the fluorescence disturbs the Raman measurement a coupled attenuated total reflection infrared (ATR-IR) unit can be used. The first example of such analyses was the determination of the depth profile of fire-induced degradation of ethylene vinyl acetate (EVA) which could hardly be determined by any other methods, such as thermogravimetric analyzer coupled with Fourier transform infrared (TG-FTIR) spectroscopy used previously to examine the pyrolysis of EVA.5–7 It was indicated that on the surface of heattreated EVA the oxidation of polyethylene (PE) segments is dominant and leads to the formation of a carbonaceous surface layer, while in the deeper region scission of the acetate side-chains in vinyl acetate segments is characteristic. Even deeper, at 40 mm, melted (but not degraded) polymer could be found. Laser treatment was found to be comparable with the effect of a cone heater. Figure 6.1 shows that the measured points of different treatments fall on the same curve: 50 seconds of cone treatment, in the given experiment, corresponds to about B1 second of laser treatment, and both represent B2.5 MJ m
2. Furthermore, by adjusting the energy of the laser the rate of degradation, according to the need, can be slower or even higher. However, these preliminary experiments indicated that additional analyses were needed to predict the fire behaviour of polymeric materials. Comprehensive characterization should consider:  heat distortion temperature [influences the reproducibility of single burn item (SBI) test];  heat conductivity–temperature function (influences ignitability and heat feedback);  viscosity–temperature function (influences gas transport, dripping, spread of fire);  chemical composition of gas phase (influences ignitability and flame propagation);  temperature distribution in the gas and solid phases (influences heat feedback);  heat balance–time function of the combustion process (influences fire persistence);  ratio and rate of the condensed phase mechanism (influence barrier-layer formation);  size and kinetic energy of bubbles (influence the uniformity of the barrier layer);

76

Chapter 6

Figure 6.1

Correlation between the degradation caused by cone heater treatment and CO2 laser treatment on EVA samples.

 chemical composition and physical characteristics of the condensed phase, including the size, structure, barrier capacity, mechanical and heat resistance of char (influence the mass and heat balance). It is quite challenging that no equipment that adequately performs such a complex analysis is available, especially for micro-scale characterization. This work attempts to make progress in the characterization by combining the LP–Raman microscope system with FTIR gas analysis and micro-thermal analyser (to determine temperature-dependent deformation and heat conductivity). The use of a rheometer at high temperature for small-scale measurements of the physical resistance of the char is published elsewhere.8,9 Concerning the mechanism of fire retardancy, a better understanding is required, especially to clarify the role of nanofillers, which is still a challenge in spite of the intensive studies. The relationship between interfacial (e.g. particle– particle and polymer–particle) interactions and fire performance has not been elucidated either. The interaction at the interfaces in flame retarded polymers can be considered from different points of view:  interactions in the solid state influence mechanical, thermal and electrical properties, and heat deflection temperature;

Complex Micro-analysis Assisted Design of Fire-Retardant Nanocomposites

77

 interactions in the melt state influence processability and melt dripping;  interfacial interactions in the course of degradation influence the condensed-phase and, indirectly, the gas-phase mechanisms. Interactions in the solid state (including covalent bonds, acid–base interactions, van der Waals interactions and anchoring-type junctions) that are formed between the phases of multicomponent systems, contact area, thickness and structure of the interfacial layers may influence the temperature-dependent mechanical properties. The dimensional stability under the effect of fire strongly influences the response of material to certain types of measurement set-up, such as SBI and Federal Motor Vehicle Standard Safety n1302 (FMVSS) tests. Increased adhesion at the interfaces improves the physical stability, but high local stresses may occur at the interfacial zone because of different heat expansion of adjacent phases. An interlayer of higher thickness, formed by an elastomer, for instance, decreases the chances of formation of local stress peaks and of consequent interfacial debonding.10 Optimally selected interfacial additives may improve both the wetting (surfactant effect) and adhesion (coupling effect) of these inclusions. Coupling agents, if introduced into the appropriate interfacial zone, build chemical bridges between the phases.11,12 Reactive surfactants (RS), considered to be targeted reactive interphase (IP) modifiers, have been developed to combine the benefits of surfactants and coupling agents.13,14 Use of RS additives has resulted in improved mechanical properties in polymer blends, while in pigmented PE it contributed to higher photostability.15–17 However, better interfacial interaction does not automatically ensure better fire retardancy. The term ‘adaptive IP’ has been suggested recently to describe interlayers responding to the changes in their environment that act as a trigger signal to initiate the desirable transformations.18–20 Heat-induced IP transformations may facilitate stabilization by controlling the degradation process. Interactions in the melt state in presence of fillers, especially nanofillers of high specific surface area, increase the viscosity of the polymers as the secondary bonds form between the surface of the inclusions and the macromolecules restricting the free displacement in the direction of the shear tensor. The viscosity-increasing effect and its influence on fire retardancy have been analyzed for intumescent systems and fire retarded polymers that contain montmorillonite (MMT).21–23 The effective contact area between the phases is influenced by the value of wetting, especially at interfaces of irregular shape. Polar components of interacting phases and acid–base interactions promote maximal wetting. The kinetics of the wetting process depend on the viscosity of the melt and on the value of the shear force. The need for high viscosity to avoid melt dripping seems to be in contradiction with the processability requirements, but a material with strongly shear-rate dependent viscosity can meet both requirements. Degradation-related interactions between the polymer matrix and the included particles involve chemical reactions at the phase borders and/or physical adsorption. The catalytic or chemical effect of the active surface of the nanofillers may lead to early mass loss, while the same effect may also promote the advantageous charring process at higher temperatures. Thus, the mechanism of

78

Chapter 6

degradation can be influenced by IP modification using catalyst atoms or other active species, with different final consequences on fire retardancy from case to case.24 Adsorption at the interfaces may lead to considerable extraction of the stabilizer(s) from the polymer matrix, which results in earlier degradation compared to the pristine matrix.16,17,25 This chapter reports work that aims to make a step towards a better understanding of the complex relationship between structure and fire performance.

6.2 Experimental 6.2.1 Materials The materials used were:  Polypropylene (PP): Moplen HP400R type (Basell Polyolefins), melt flow index 25 (at 230 1C/2.16 kg).  Maleated polypropylene (PP-g-MA): Licomont AR 405 type (Clariant GmbH), acid content 37–45 (mg KOH g
1), density 0.89–0.93 g cm
3 at 23 1C.  EVA copolymer: IBUCELL K 100 with vinyl acetate content of 28% (H.B. Fuller).  Multiwall carbon nanotube: MWCNT Nanocyl 3100 type (Nanocyl S.A., Belgium), average diameter 10 nm, length 0.1–10 mm, metal oxide content o5%.  Sepiolite: Pangel S9 type (Tolsa S.A., Spain), other mineral content: 15%.  Na-MMT: Microtec type (Eurotrade Ltd., Hungary), mass loss 7.48% at 10001C. The components of the P-epoxy resin interlayer were:  Eporezit AH-16 (non-modified, resin-like reactive dilutant, epoxy equivalent 160–175; viscosity 800–1800 mPa.s at 25 1C; density 1.24 g cm
3 at 25 1C:, hydrolyzable chlorine content 1.5 mass per cent) supplied by P+M Polimer Ke´mia Kft., Hungary  TEDAP a newly synthesized phosphorus-containing reactive amine (amine number 510–530 mg KOH g
1; viscosity 400 mPa.s at 20 1C; curing time 7 days at 25 1C, 4 hours at 80 1C).26 The ratios of components in PP : PP-g-MA : P-epoxy : nanofiller systems were 60 : 20 : 15 : 5; for the carbon nanotube (CNT) the ratio was 63 : 21 : 15.8 : 0.2 because a higher ratio of CNT could not be perfectly homogenized in epoxy resin.

6.2.2 Methods Samples were prepared in a Brabender PL 2000 with a rotor speed 50 1 min
1 at 180 1C, the duration of compounding being 10 minutes.

Complex Micro-analysis Assisted Design of Fire-Retardant Nanocomposites

Figure 6.2

79

Set-up of laser pyrolysis and analysis of solid and gas phases.

The LP system (see Figure 6.2) comprises a CO2 laser (SYNRAD 48-1) and a Raman-IR microscope (from HORIBA Jobin Yvon Inc.). The beam of the CO2 laser on the sample falls within the IR range so its effect on the sample simulates the influence of fire. The LP unit is equipped with a Universal Laser Controller, which allows the operator to control the laser power using pulse width modulation (PWM). The laser was used at the standard 5 kHz frequency and the laser power was controlled. Gate control provides the ability to vary laserexposure period. The pyrolysis unit is mounted on a flange and the CO2 laser beam is focused on the sample surface using a special mirror. Before laser exposure, the sample is placed on the top of a probe, which is inserted into a position just below the Raman–ATR microscope system. The gases that evolved during the laser irradiation were evacuated and carried to IR gas cell for further analyses. The subsequent chemical changes at the surface are then monitored, either by Raman microscopy or via an ATR technique. Raman analyses to characterize the structure and interactions in the solid state were performed using a LabRam-type confocal Raman microscope (Horiba Jobin Yvon, France). The magnification used was 50 during the measurement and the excitation source was a frequency-doubled Nd–YAG laser emitting at 532 nm. A microthermal analyser of the m-TA 2990 type (TA Instruments) was used to determine the microthermal deformation and thermal conductivity values. The instrument is an AFM system into which a controllable wire is inserted, as a tip, in the AFM cantilever and used to scan along the surface of the sample (Figure 6.2). The heating rate at deformation was 5 1C s
1 in the temperature range 25–250 1C, the contact during the heating was set with the Z-piezo to the value of 20 V. The heat conductivity correlates with the power (F in mW) necessary to maintain the temperature of the probe, which was chosen to be 100 1C. The relative heat conductivity of the samples is expressed as the dimensionless value l/lair ¼ Fsample (mW)/Fair (mW), where Fair ¼ 1.01 mW. Scanning electron microscopy (SEM) analysis of the additives was made using a JEOL 5500 LV instrument.

80

Chapter 6

Heat treatment of the samples in a rheometer (before residue analysis) was performed using an AR 2000 system from TA Instruments Co. The fire performance was determined using the horizontal burning test (ISO 3795 standard) to establish the rate of flame spreading (burning) and dripping. Also investigated were the limiting oxygen index (LOI), according to ASTMD 2863, and mass loss calorimetry (Fire Testing Technology) according to ISO 13927 with heat flux 50 kW m
2, sample surface area 100 cm2, Al tray of 10  10  0.3 cm size, three replicate samples (covariance 3%), exhaust gas flow 0.024 m3 s
1.

6.3 Results and Discussion To demonstrate the applicability of the small-scale modelling and micro-analytical approach, described in the experimental part, preliminary results are presented on consecutive steps of development of a new fire retardant–PP system.

6.3.1 Nanonetwork Formation To study the nano-effects in the condensed and gas phases, PP–MWCNT systems were analyzed. In this case the degradation initiated by the laser treatment created a hole, the size of which was influenced by the duration of the treatment and the composition. Table 6.1 shows that the time-dependent growth of the hole is less when nanofillers, especially CNT, are present. Obviously, nanocomposites with a higher heat conductivity resulted in a lessfocused heating effect and thus reduced degradation in the centre of the treatment. For the PP–MWCNT system the structure of the hole was analyzed by SEM. A fine CNT network was found within the hole, which was uniform in each hole analyzed (see Figure 6.3). The SEM image confirms our earlier (Raman) Table 6.1

Depth of the hole induced by treatment (of increasing duration) with CO2 laser on the surface of different PP systems. d Sample

Time Sample

0.2 sec

0.4 sec

0.6 sec

0.8 sec

1 sec

2 sec

5 sec

PP-MWNT 3% 128 mm 356 mm 489 mm 635 mm 726 mm 1856 mm 2330 mm PP-sepiolite 3% 257 mm 532 mm 1166 mm 1747 mm 2021 mm 3825 mm 44000 mm PP 413 mm 966 mm 1233 mm 2816 mm 3752 mm 44000 mm 44000 mm

Complex Micro-analysis Assisted Design of Fire-Retardant Nanocomposites

Figure 6.3

81

Internal structure of the hole induced on the surface of a PP–3%MWCNT system by 5 second CO2 laser treatment.

observation that the residue after combustion of PP–MWCNT system consists of pure CNT.27 The network structure was expected to increase the dimensional stability and heat conductivity. For a micro-scale comparison of the heat distortion and the heat conductivity of the PP and PP–MWCNT systems the m-TA method was used. The results given in Figure 6.4 demonstrate the effect of the CNT network on the temperature-dependent deformation of the PP. Up to B150 1C the thermal dilation of the samples can be seen, and is considerably reduced in the presence of CNT. Forced penetration of the AFM tip into the materials starts when the compliance, determined by the mobility of the macromolecules, becomes high enough. In the presence of CNT the mobility of macromolecules at the interfacial region is restricted by the interfacial interaction; therefore, the deformation starts at higher temperature and takes place in a less steep manner than in the case of pristine PP. This result suggests there is less chance of distortion and dripping because of fire if CNT is dispersed in the polymer matrix. The heat conductivity can be determined by keeping the tip of the m-TA system at a constant elevated temperature. The energy needed to maintain this temperature when the tip is in contact with the material is proportional to its heat conductivity. This method allows the determination of micro-scale local differences up to the temperature of combustion, which cannot be done with other heat-conductivity measurements. The first experiments confirmed the applicability of m-TA to characterize the studied samples, as the results given in Table 6.2 correlate well with recently published heat-conductivity data of PP– CNT samples.19

82

Figure 6.4

Chapter 6

The temperature-dependent deformation of PP and PP–3%MWCNT determined by the m-TA method.

Even these preliminary results clearly show the considerable increase in heat conductivity with CNT in the composites, and its opposite effect on the char residue. The introduction of 3% MWCNT into PP increased this value to thrice that of the reference PP, which suggests a good contact between the polymer matrix and the CNT of high heat conductivity. The opposite tendency was found when the heat conductivity of the residue was measured. The CNT residue (due to its voluminous, loose fibrous structure) isolates the heat 10 times better than does the char residue of an EVA (which was taken as reference in this case, as for PP no residue remains to be measured). The effects of changes to the heat conductivity on fire retardancy may be ambiguous. The increased heat conductivity of PP–CNT reduces the time to ignition (TTI); however, if it forms a heatresistant thermal barrier on the surface of material, this will slow down the degradation in the subsequent stage of combustion. This effect of increased heat conductivity explains why the CNT decreases the TTI, as reported earlier, even though the degradation of PP is shifted to higher temperature (moderately or considerably under nitrogen or air atmosphere, respectively) in the presence of CNT.27,28 As the combustion process continues, the heat barrier effect of the char, which consists of CNT residue of low heat conductivity, results in a reduced heat release rate (HRR).28 However, according to our experience the char is not consistent enough to resist the intensive bubbling in the later stage of cone heater treatment at 50 kW m
2. The photo of the residue in Figure 6.5 shows the bubbling-induced removal of the CNT from the central part of the material. The decomposition of CNT is excluded as its amount in the residue was equal to that of the introduced CNT. The entanglement of nanofibres can be preserved only if mild flame treatment conditions are applied or if binding sites are introduced.24,29

83

Complex Micro-analysis Assisted Design of Fire-Retardant Nanocomposites

Figure 6.5

Photograph of the residue of PP–3%MWCNT after treatment under cone heater at 50 kW m
2.

Table 6.2

Conductivity data and results of gas phase analysis of PP and PP + CNT systems.

PP PP + 3%MWCNT EVA (reference)

Relative heat conductivity l/lair

Integrated values of gas components in gas phase [a.u.]

reference

residue

CO2

CO

Org.

1.06 1.19 1.07


1.02 1.25

154600 148500

37.2 35.6

8.1 5.0

Note: Org. ¼ organic components; a.u. ¼ arbitrary unit

The indirect effect of the described condensed-phase changes in the gas phase was analyzed by the FTIR unit coupled to the laser pyrolyser. A moderate decrease in all the decomposition gases was found as a consequence of the transport-restriction effect of the CNT network (see Table 6.2). This moderate effect is not enough to provide real fire retardancy: the LOI value of 19 was found independantly of the composition for PP–MWCNT composites with 0.5, 1, 2, 3, 4, and 8% CNT content while the burning rate decreased and no dripping occurred above 2% MWCNT (Figure 6.6).

84

Figure 6.6

Chapter 6

Rate of horizontal flame spread and dripping against the percentage of MWCNT in PP.

The PP–CNT system demonstrated the effects of inter-particle interactions, among which the maintenance of percolated structure seems to be especially important; however, the introduction of CNT could not provide fire retardancy at an industrially important level. Based on the previous results, which indicate the role of interface modification in fire retardancy, we have attempted to produce a polymer layer of fire retardant activity around nanofillers.30

6.3.2 Intumescent Polymeric Particle Formation Epoxy resin was selected as a polymeric interlayer around different nanofillers, i.e. CNT, MMT and sepiolite (fibrous clay). The selection was based on previous results.26,31 A new cross-linking agent (TEDAP), which is a reactive fire retardant, has been synthesized recently for epoxy resins.26 Detailed mechanistic studies made clear that its excellent performance is established by an optimal balance and timing of the gas-phase and solid-phase mechanisms.32 Early gas-phase action may compensate the TTI-reducing effect of CNT and clay. The laser treatment of this inherently flame retardant epoxy type resulted in a foamed structure, rather than the hole formed in PP, accompanied by an intense charring process (see Figure 6.7). Based on the considerable intumescent effect of the new fire retardant epoxy the idea to use it in combination with nanoparticles as fire retardant additives in PP emerged.

Complex Micro-analysis Assisted Design of Fire-Retardant Nanocomposites

85

Figure 6.7

Surface image of P-epoxy resin after 1 second treatment with CO2 laser.

Figure 6.8

Preparation method for forming PP–P-epoxy–nanofiller system.

The preparation method for the samples is described in Figure 6.8. The nanoparticles were dispersed in epoxy monomers (the high polarity of epoxy resin facilitated good dispersion) and in the course of the curing process it was mixed with a PP–MA-g-PP matrix. Surprisingly, the formed structure, shown in Figure 6.9, was the same in all cases and independent of the geometry of the nanofiller. Very small droplets were formed in a good dispersion. It is assumed that the droplet form is preferred because of the high surface energy of the epoxy resin, and the nanofillers occur within the droplets in folded and/or lamellar form. The PP nanocomposites were heat-treated in a rheometer. Here only the analysis of the residue is presented, while the rheological results will be published elsewhere. Micro-Raman comparison of the residues of PP loaded with pure MWCNT and of the system containing epoxy-coated MWCNT demonstrates the significant difference between the two materials. Figure 6.10

86

Figure 6.9

Chapter 6

SEM image of the surface of a fracture of a PP–P-epoxy–sepiolite sample broken under liquid N2.

shows that, in the case of PP–MWCNT, the residue is pure MWCNT, which, according to the ratio of D (B1340 cm
1) and G bands (1500–1600 cm
1), is even purer than the original CNT. Thus in this case the combustion acts as a purification process. The presence of an epoxy resin interlayer dominates the structure of the residue, which shows the wide bands of amorphous carbon, indicating the transformation of a considerable amount of polymer into amorphous char. The PP systems that contain nanofillers coated with epoxy resin were compared by mass loss calorimetry. It is interesting to see in Figure 6.11 the slightly increased TTI of the composites that contain these coated nanofillers. This improvement is noteworthy when the decreasing effect of all the nanofillers studied is considered. It confirms the importance of a gas-phase fire retardant mechanism (of fire retardant–epoxy-resin units) in the initial stage of fireinduced degradation. The largest reduction of HRR and the most consistent char was achieved, when the active epoxy-resin interlayer was applied on MMT. Our assumption for explaining these results considers the expansion capacity of the fire retardant–epoxy-resin droplets that contain the nanofillers. The fibrous inclusions may restrict the expansion, while the MMT influences the foaming process of epoxy resin advantageously, as reported previously.33 Expansion of the tiny particulates, shown in Figure 6.9, may reach the percolation level. In conventional intumescent systems a continuous char layer is formed only after the decomposition of a considerable amount of matrix polymer. In contrast, in a ‘percolating intumescent system’ the expanding particles reach each other in the early stages of the process, which provides

Complex Micro-analysis Assisted Design of Fire-Retardant Nanocomposites

87

Figure 6.10

Raman spectra of PP–MWCNT systems, their residue and reference materials.

Figure 6.11

HRR curves of PP, PP–P-epoxy, PP–P-epoxy–sepiolite; PP–P-epoxy– MMT samples (external heat flux 50 kW m
2).

prompt protection to the underlying polymer phase.34 (A percolation model for isodimensional inclusions of polystyrene blend has been applied recently.35) Parameters that influence the percolation are the interparticle distance, volume increase and rate of action. Based on the results reported in this chapter, it is

88

Figure 6.12

Chapter 6

Scheme of percolation of intumescent particles and nanofibres.

assumed that nanotubes and/or nanofibres in the PP phase that interconnect the intumescent particles, according to Figure 6.12, would provide an even better performance. Validation of this hypothesis requires further investigation.

6.4 Conclusion The complexity of the features to be considered when fire retardancy is designed involves the need to develop complex structures, small-scale modelling and coupled analytical techniques. Treatment with a CO2 laser is an accurate tool with which to model the fire effect, while the set of micro-Raman–ATR, FTIR, mTA and HT-rheology measurements seem to be suitable for providing complex, rapid information about the changes in the gas and solid phases. Micro-scale determination of the time function to thermal balance of the combustion process requires further development. Results on PP–nanofiller systems confirm that their main fire retardancy advantages are the improved form stability at high temperatures and decreased dripping. Introducing P-containing a fire retardant, thick epoxy resin layer around nanofillers increases the TTI and reduces the HRR. This is ascribed to a balanced ratio of gas- and condensed-phase mechanisms and percolation of

Complex Micro-analysis Assisted Design of Fire-Retardant Nanocomposites

89

intumescent inclusions at the early stage of combustion, but these assumptions need to be validated by further analyses. The mechanistic conclusions in connection with the interfacial phenomena are:  Interfacial interaction between the nanofillers allows them to act as a physical network in the polymer matrix and so decrease the chance of deflection of the material at high temperatures and thus promote better performance in larger scale vertical fire tests.  This interaction and the formation of immobilized macromolecular interlayers around nanofillers give rise to the observed anti-dripping effect and consequently reduced spread of fire. Optimal balance between this effect and the processability requires melt of high shear sensitivity.  Fire retardant (intumescent) interfacial layers may shift the TTI and reduce the rate of heat release. The larger the expansion of the interlayer the greater the chance that rapid percolation of the intumescent particles will make the fire protection more efficient.

Acknowledgements The authors acknowledge the financial supports received through the EU-6 Framework Program (NMP3-CT-2004-505637), Multihybrids (IP 026685-2), Hungarian Research Found OTKA T049121 and Found of European Union and Hungarian state GVOP/3.1.1.-2004-0531/3.0.

References 1. D. Price, G.J. Milnes, P.J. Tayler, J.H. Scrivens and T.G. Blease, Polym. Degrad. Stab., 1989, 25(2-4), 307. 2. D. Price, G.J. Milnes, C. Lukas and T.R. Hull, Int. J. Mass.Spectrom., 1984, 60, 225. 3. D. Price, G. Fengge, G.J. Milnes, B. Eling, C.I. Lindsay and P.T. Mcgrail, Polym. Degrad. Stab., 1999, 64(3), 403. + G. Marosi, Period. Polytech. Chem., 2008 (accepted). 4. B.B Marosfoi, 5. A. Marcilla, A. Go´mez and S. Menargues, J. Anal. Appl. Pyrol., 2005, 74(1–2), 224. 6. B. Marosfoi, A. Szabo´, A. Toldy, P. Anna, Gy. Marosi, D. Tabuani and G. Camino, in Recent Advances in Flame Retardancy of Polymeric Materials XVIII, ed. M. Lewin, BCC Inc., Norwalk USA, 2006, pp. 25. 7. B.B. Marosfoi, G. Marosi, A. Szabo, B. Vajna and A. Szep, Polym. Degrad. Stab., 2007, 92(12), 2231. 8. M. Bugajny, M. Le Bras and S. Bourbigot, Fire Mater., 1999, 23(1), 49. 9. P. Anna, Gy. Marosi, I. Csontos, S. Bourbigot, M. Le Bras and R. Delobel, Polym. Degrad. Stab., 2001, 74(3), 423.

90

Chapter 6

10. Gy. Marosi, Gy. Bertalan, I. Ruszna´k and P. Anna, Colloid. Surface., 1986, 23(3), 185. 11. W. Gutowski, S. Li, L. Russell, C. Filippou, P. Hoobin and S. Petinakis, Compos. Interface., 2002, 9(1), 89. 12. W.S. Gutowski, J Adhesion, 2003, 79(5), 445. 13. Hungarian Patent 218 016 (2000). 14. Gy. Marosi, P. Anna, I. Csontos, A. Ma´rton and Gy. Bertalan, Macromol. Symp., 2001, 176, 189. 15. G. Bertalan, G. Marosi, P. Anna, I. Ravadits, I. Csontos and A. To´th, Solid State Ionics, 2001, 141–142, 211. 16. M.A. Maatoug, P. Anna, Gy. Bertalan, I. Ravadits, Gy. Marosi, I. Csontos, A. Ma´rton and A. To´th, Macromol Mater. Eng., 2000, 282, 30. 17. P. Anna, Gy. Bertalan, Gy. Marosi, I. Ravadits and M.A. Maatoug, Polym. Degrad. Stab., 2001, 73(3), 463. 18. Gy. Marosi, Gy. Bertalan, in Modification and Blending of Synthetic and natural Macromolecules, eds. F. Ciardelli, S. Penczek, NATO Science Series, Kluwer Acad. Publ. Dordrecht, 2004, 175, 1351. 19. I. Luzinov, S. Minko and V.V. Tsukruk, Progr. Polym. Sci., 2004, 29(7), 635. 20. S. Keszei, Sz. Matko´, Gy. Bertalan, P. Anna, Gy. Marosi and A. To´th, Eur. Polym. J., 2005, 41, 697. 21. S. Bourbigot and M. Le Bras, in Fire Retardancy of Polymers: The Use of Intumescence, eds. M. Les Bras, G. Camino, S. Bourbigot, R. Delobel, The Royal Society of Chemistry, Cambridge, 1998, pp. 222. 22. P. Anna, Gy. Marosi, S. Bourbigot, M. Le Bras and R. Delobel, Polym. Degrad. Stab., 2002, 77(2), 243. 23. P. Anna, Gy. Marosi, Gy. Bertalan, A. Ma´rton and A. Sze´p, J. Macromol. Sci. Phys., 2002, B41(4–6), 1321. + B. Bodzay, Gy. Marosi, J. Therm. Anal. 24. A. Szabo´, B.B. Marosfoi, Calorim., (submitted). 25. B. Mailhot, S. Morlat, J. Gardette, S. Boucard and J. Duchet J. Ge´rard, Polym. Degrad. Stab., 2003, 82(2), 163. 26. A. Toldy, N. To´th, P. Anna and Gy. Marosi, Polym. Degrad. Stab., 2006, 91(3), 585. + A. Szabo´, Gy. Marosi, D. Tabuani, G. Camino and 27. B.B. Marosfoi, S. Pagliari, J. Thermal Anal. Calorim., 2006, 86(3), 669. 28. B. Schartel, P. Potschke, U. Knoll and M. Abdel-Goad, Eur. Polym. J., 2005, 41, 1061. 29. T. Kashiwagi, E. Grulke, J. Hilding, K. Groth, R. Harris, K. Butler, J. Shields, S. Kharchenko and J. Douglas, Polymer, 2004, 45(12), 4227. 30. Gy. Marosi, P. Anna, A. Ma´rton, Gy. Bertalan, A. Bo´ta, A. To´th, M. Mohai and I. Ra´cz, Polym. Adv. Technol., 2002, 13(10–12), 1103. 31. N. Abacha, M. Kubouchi, K. Tsuda and T. Sakai, Exp. Polym. Lett., 2007, 1(6), 364.

Complex Micro-analysis Assisted Design of Fire-Retardant Nanocomposites

91

32. A. Toldy, A. Szabo´, Cs. Nova´k, J. Madara´sz, Gy. Marosi, Polym. Degrad. Stab., (doi:10.1016/j.polymdegstab.2008.02.011). 33. A. Toldy, P. Anna, I. Csontos, A. Szabo, Gy. Marosi, Polym. Degrad. Stab., 2007, 92(12), 2223. 34. F. Zhang, J. Zhang and Y. Wang, Exp. Polym. Lett., 2007, 1(3), 157. 35. Z. Guo, Z. Fang and L. Tong, Exp. Polym. Lett., 2007, 1(1), 37.

Nanoparticulate Fillers

CHAPTER 7

Impact of Nanoparticle Shape on the Flammability of Nanocomposites F. YANG, I. BOGDANOVA AND G. L. NELSON College of Science, Florida Institute of Technology, Melbourne, Fl, 32901, USA

7.1 Introduction There are a number of advantages of polymer–inorganic nanocomposites when compared to conventional composite materials.1–5 The surface area of nanoparticles is significantly larger than that of micro-sized particles, which provides a larger surface contact for reinforcement of polymers. The amount of nanosized additives needed is much less than those of conventional fillers to achieve the same level of reinforcement. Although polymer–inorganic nanocomposites have been studied for decades, there is limited understanding of their reinforcement and degradation mechanisms. The problems associated with controlling factors that affect the reinforcement and degradation of polymer nanocomposites are not resolved. In previous studies,6–8 the dispersion (or distribution) of nanoparticles in polymer nanocomposites was considered to be the most important factor in producing strong interfacial interactions. However, even good dispersion of nanoparticles, as achieved by many researchers, does not necessarily improve the physical properties of nanocomposites. This suggests that factors other than good dispersion of nanoparticles determine the enhancement of properties for the final composite.

Fire Retardancy of Polymers: New Strategies and Mechanisms Edited by T Richard Hull and Baljinder K Kandola r Royal Society of Chemistry 2009 Published by the Royal Society of Chemistry, www.rsc.org

95

96

Chapter 7

To understand the mechanism of thermal degradation and flammability, factors such as the nature of polymer matrices and nano-additives, and the effect of shape, size and loading of the nano-additives on interfacial interactions in polymer nanocomposites are investigated.

7.2 Experiment 7.2.1 Preparation of Polymer–Inorganic Nanocomposites Styrons 685D polystyrene (PS) for the preparation of PS–Aerosils 90 nanocomposites, Styrons 685P for PS–Aerosils R972 nanocomposites, Styrons 685DW for PS–alumina nanocomposites, developmental grade polycarbonate resin were obtained from Dow Plastics Inc. Lexans 103 and GEs OQ 3620 110 polycarbonate (PC) were provided by GE Plastics, Inc. Nanoscale silica, Aerosils 90 (20 nm, spherical) and Aerosils R972 (16 nm, spherical – an organosilica where the surface is covered by methyl groups) were provided by Degussa Co. Aerosils 90 was pretreated with silanol-terminated polydimethylsiloxane (PDMS) and phenyltrichlorosilane (PTCS) for PS–silica nanocomposites. Aerosils R972 was used as received. Alumina hydrate (nanorods, 100 nm/10 nm, Aluminasols 200) in the form of methanol–water colloidal solution was purchased from Nissan Chemical Inc., Japan. Alumina, alpha form (40–80 nm, irregular) and alumina hydroxide (15 nm, spherical) were purchased from Nanostructured and Amorphous Materials Co., USA. Alumina whiskers (2800 nm/2–4 nm) were purchased from Sigma-Aldrich, USA. 3-Methacryloxypropyltrimethoxysilane (MAP) and methyltrimethoxysilane (MT) were purchased from Gelest Inc., USA. Polymer– inorganic nanocomposites were prepared by a single-screw extrusion method developed in the laboratory using a three-quarter inch CW Brabender Table Top Independent Extruder.5

7.2.2 Mechanical Testing of Polymer–Inorganic Nanocomposites A Tinius Olsen Series 1000UTM tensile tester with an analogue recorder was used for tensile strength, modulus and elongation at break for all materials, according to ASTM D638. The testing rate was 0.05 inch min
1.

7.2.3 Morphology Study for Polymer–Inorganic Nanocomposites The samples for transmission electron microscopy (TEM) were prepared by a Leica Ultracut UC6 Microtome. Glass knives (prepared with LKB 7800 Knifemaker) and a diamond knife (Electron Microscopy Sciences, USA) were used. The thickness of the sample is 150 nm and formvar-coated copper grids with 100 mesh (Electron Microscopy Sciences, USA) were used to support the nanocomposites sections for TEM. Tests were conducted using a Zeiss EM900 transmission electron microscope equipped with a Morada Soft Imaging System digital CCD camera, Olympus, USA.

97

Impact of Nanoparticle Shape on the Flammability of Nanocomposites

7.2.4 Thermal Degradation of Polymer–Inorganic Nanocomposites Thermal gravimetric analysis (TGA) for polymer–inorganic nanocomposites was performed using a Hi-Res TGA 2950 thermogravimetric analyzer (TA Instruments, USA). The polymer–alumina samples were pre-heated to 100 1C, equilibrated at 100 1C for 3 minutes to eliminate moisture and solvent, and then heated from 100 1C to 650 1C at 20 1C min
1. The polymer–silica samples were pre-heated to 100 1C and held for 5 minutes to eliminate solvent and moisture, followed by heating to 550 1C at 10 1C min
1. All tests were run under helium atmosphere.

7.2.5 Flammability of Polymer–Inorganic Nanocomposites Flammability of nanocomposites was tested by a horizontal burning test according to ASTM D635, vertical burning test according to ASTM D3801, oxygen index (LOI) according to ASTM D2863 and cone calorimetry. Cone calorimetry studies were performed on a Custom Scientific Cone Calorimeter instrument at 35 kW m
2, according to ASTM 1354. A modified cone-calorimetry procedure was used. Samples were made of two conjunctive strips (5  10 cm) with 1.6 mm thickness instead of 10  10  2.5 cm samples. Five tests were conducted for each sample (ASTM 1354). The error in the results is less than 5%.

7.3 Results and Discussion 7.3.1 Polycarbonate–Inorganic Nanocomposites 7.3.1.1

Morphology Studies

The TEM image of PC–alumina 40–80 nm modified with MAP is shown in Figure 7.1(a). Modified alumina nanoparticles were dispersed in PC and appeared spherical in shape. Figure 7.1(b) shows a TEM image of PC–alumina

a

3820

Figure 7.1

b

1000nm

63700 : 1

500nm

TEM of polymer/alumina a) 5 wt% PC/alumina 40–80 nm modified with MAP, b) 5 wt.% PC/alumina 100 nm modified with MAP.

98

Chapter 7

100 nm/10 nm nanoparticles modified with MAP. Alumina nanoparticles (appear as dark needles) were well dispersed in PC. As indicated from Figures 7.1(a) and 7.1(b) nanoparticles were well dispersed in the PC matrix.

7.3.1.2

Thermal Stability

Degradation mechanisms for polymer–inorganic nanocomposites are based on studies of thermal degradation and flammability, which include the effect of nanoparticles on the degradation pathway for the polymer. As reported,9 the chemical nature of nanoparticles and polymers affects the chemical degradation processes in nanocomposites. Several factors affect the thermal degradation of PC–alumina nanocomposites. It was observed10,11 that aluminium oxide can catalyze the decomposition of carbonyl-containing polymers, which results in carboxylic acid elimination. Moreover, the degradation products, free radicals, are more reactive on the alumina oxide surface. PC has main-chain carbonyl groups, which can be eliminated in heterogeneous catalysis by alumina nanoparticles and result in carboxylic acid. The thermal degradation of PC–alumina nanocomposites is expected to proceed faster as a result of alumina-catalyzed PC decomposition. The elimination of carboxylic acid from PC leads to a decrease in the PC chain length. Thermal degradation mechanisms were studied using TGA and scanning electron spectroscopy (SEM) on the chars of PC– alumina and PC–silica nanocomposites. Thermal stability is studied on the basis of temperatures at 10% and 50% weight loss. The results of TGA for PC– alumina nanocomposites are listed in the Table 7.1. PC–alumina nanocomposites with 1–3% of alumina hydrate whiskers showed a 20–21 1C increase at 10% weight loss for PC–1% alumina whisker nanocomposite. Alpha-alumina (40–80 nm, spherical) exhibited decreased thermal stability for 5 weight per cent (wt%) PC–alumina nanocomposites. On the other hand, PC–silica nanocomposites exhibit a different effect of the nanoparticles on the thermal degradation process, as is evident from Table 7.212. Silica affects the thermal degradation pathway by a restriction mechanism, which results in enhanced char formation for char-forming polymers. The higher loading of silica in PC–silica nanocomposites resulted in a larger number of restriction sites for PC chain mobility, which enhanced char formation and the thermal stability of PC–silica nanocomposites compared to those of PC. For instance, a 15 1C increase at 50% weight loss for PC–3% silica is observed. Another factor that affects the thermal stability of PC–silica nanocomposites is the trapping of degradation products, radicals, on the silica surface, which has a stabilizing effect on the thermal degradation process for PC–silica nanocomposites. To gain a better understanding of the PC–alumina and PC–silica degradation processes, SEM images of thermal degradation residues for PC–alumina and PC–silica nanocomposites were taken. Residues were prepared by heating samples to 400 1C, 460 1C and 500 1C in a helium atmosphere. Images are shown in Figure 7.2. As clearly indicated from Figure 7.2(a), alumina particles are combined in larger aggregations on the

99

Impact of Nanoparticle Shape on the Flammability of Nanocomposites

TGA of PC–alumina nanocomposites.

Table 7.1

10% loss

50% loss

Sample

T (1C)

DT (1C)

T (1C)

DT (1C)

PC (Dow, developmental resin) PC-Al-100-MAP-0.5 PC-Al-100-MAP-1 PC-Al-100-MAP-2 PC-Al-100-MAP-5 PC-Al-40-MAP-0.5 PC-Al-40-MAP-1a PC-Al-40-MAP-2 PC-Al-40-MAP-5 PC (GE OQ 3620) PC-Al-15-MAP-0.5 PC-Al-15-MAP-1 PC-Al-15-MAP-2 PC-Al-15-MAP-5 PC-Al-wh-MAP-1 PC-Al-wh-MAP-3 PC-Al-wh-MAP-5 PC-Al-wh-MT-1 PC-Al-wh-MT-3 PC-Al-wh-MT-5

460 460 461 472 465 462 469 468 435 463 481 480 485 478 475 483 479 484 477 477

0 0 1 12 5 2 9 8
25 0 18 17 22 15 12 20 16 21 14 14

501 500 505 509 508 503 508 508 497 505 512 511 510 512 507 513 510 513 511 511

0
1 4 9 8 3 8 8
3 0 7 6 5 7 2 8 5 8 6 6

a

Example of sample code structure: in PC-Al-40-MAP-1, PC abbreviates polycarbonate matrix, Al means alumina nano-additive, 40 stands for the size of nanoparticles, MAP is the silane coupling agent used, 1 is the wt% of nano-additive. Al-wh is whisker-shaped alumina nano-additive.

TGA for PC–silica nanocomposites.12

Table 7.2

10% weight loss Sample s

PC (Lexan PC-Si-16-1 PC-Si-16-3 PC-Si-16-5

a

Figure 7.2

103)

50% weight loss

T (1C)

DT (1C)

T (1C)

DT (1C)

452 456 466 456

0 4 14 4

482 492 497 491

0 10 15 9

b

SEM of PC/inorganic chars, a) PC/alumina (40–80 nm, irregular), b) PC/ silica (R972).

100

Chapter 7

surface of the materials, which indicates that the alumina particles are combined on the surface. Alumina forms a layer on the surface of nanocomposites. In contrast, examination of the residues of PC–silica samples [Figure 7.2(b)] clearly indicates char formation. No silica particles were observed on the surface of degraded PC–silica nanocomposites, which indicates that silica stabilizes PC–silica nanocomposites by staying in the body of the material.

7.3.1.3

Flammability Studies

Flammability tests for PC–inorganic nanocomposites were done to investigate the input of the degradation mechanisms on the flammability of PC–inorganic nanocomposites. Cone calorimetry results for PC–alumina nanocomposites from alumina hydrate (100 nm/10 nm, needles) and alpha-alumina (40 nm, spheres) are presented in Tables 7.3–7.6. As shown in Tables 7.3–7.6, peak heat release rate (PHRR) decreased by 30% to 45% for PC–alumina nanocomposites compared to PC, which resulted from the formation of a ‘‘barrier’’ layer of alumina on the PC–alumina nanocomposite surface. As shown on the SEM images of chars, alumina Table 7.3

Cone calorimetry of PC–MT-modified alpha-alumina (40–80 nm) nanocomposites at 35 kW m
2. PC-40PC (Dow) MT-0.5

Material PHRR (kW m
2) Time to sustained ignition (s) Time of PHRR (s) THR (MJ m
2 g–1) Average mass loss rate (g m
2 s) Sample left (%)

Table 7.4

PC-40MT-1

PC-40MT-2

PC-40MT-5

1049 111

875 (–17%) 872 (–17%) 714 (–32%) 771.6 (–26) 107 46 76 67

143 1.79 33.3

132 1.76 32.9

83 1.77 26.8

104 1.72 27.8

100 1.77 26.2

24.0

25.5

22.1

22.6

27.9

Cone calorimetry of PC– MAP-modified alpha-alumina (40–80 nm) nanocomposites at 35 kW m
2.

Material PHRR, (kW m
2) Time to sustained ignition (s) Time of PHRR (s) THR (MJ m
2 g–1) Average mass loss rate (g m
2 s
1) Sample left (%)

PC-40PC (Dow) MAP-0.5

PC-40MAP-1

PC-40MAP-2

PC-40MAP-5

1049 111

950 (–9%) 80

807 (–23%) 689 (–34%) 732 (–30%) 75 83 39

143 1.79 33.3

107 1.73 32.1

108 1.82 26.0

112 1.68 24.8

112 1.86 26.8

24.0

18.8

21.6

20.3

24.2

101

Impact of Nanoparticle Shape on the Flammability of Nanocomposites

Table 7.5

Cone calorimetry of PC–MT-modified alumina hydrate 100 nm/ 10 nm nanocomposites at 35 kW m
2. PC-100PC (Dow) MT-0.5

Material PHRR (kW m
2) Time to sustained ignition (s) Time of PHRR (s) THR (MJ m
2 g
1) Average mass loss rate (g m
2 s) Sample left (%)

Table 7.6

PC-100MT-2

PC-100MT-5

1049 111

875 (–17%) 729 (–31%) 565 (–46%) 646 (–38%) 107 89 83 47

143 1.79 33.3

131 1.76 32.9

128 1.77 26.9

123 1.73 23.1

91 1.76 23.0

24

25.5

21.2

24.0

22.5

Cone calorimetry of PC–MAP-modified alumina hydrate 100 nm/ 10 nm nanocomposites at 35 kW m
2. PC-100PC (Dow) MAP-0.5

Material PHRR (kW m
2) Time to sustained ignition (s) Time of PHRR (s) THR (MJ m
2 g
1) Average mass loss rate, (g m
2 s
1) Sample left (%)

Table 7.7

PC-100MT-1

PC-100MAP-1

PC-100MAP-2

PC-100MAP-5

1049 111

802 (–24%) 676 (–36%) 580 (–45%) 645 (–39%) 76 71 67 72

143 1.79 33.3

112 1.78 28

110 1.81 25.0

109 1.58 10.2

115 1.67 22.9

24.0

17.5

18.9

21.0

23.1

Cone calorimetry of PC–silica nanocomposites.13

Material

PC (Lexan 103)

PC-Si-16-1

PC-Si-16-3

PHRR (kW m
2) Time of PHRR (s) THR (MJ m
2 g
1) Average mass loss rate (10% to 90%) Sample left (%)

940 105 2.24 27.1 17.0

900(
4.3%) 138 1.72 27.5 23.0

803(
4.6%) 117 1.65 12.4 24.7

agglomerated on the surface of degraded PC–alumina nanocomposite residues, forming a ‘‘protective barrier’’ layer. The alumina barrier layer slows the evolution of volatile gases into the burning area and protects the bulk material from the outside heat. Notice also that reduction of PHRR is 80–100 units larger for alumina hydrate nanorods than for irregular-shaped alpha-alumina– PC nanocomposites. In contrast, PHRR decreased modestly for PC–silica nanocomposites compared to PC, because of the char-formation mechanisms, as shown in Table 7.7.13 PC–alumina nanocomposites showed shorter times to ignition than PC.

102

Chapter 7

As discussed earlier, silica formed restriction sites for PC chain mobility. Char, formed as a result of PC degradation, served as an insulator for mass and energy transfer. Not only the ‘‘barrier’’ effect, but also the catalytic effect of alumina on PC degradation is involved in the flammability mechanism for PC– alumina nanocomposites. The amount of char left is less for PC–alumina nanocomposites than for PC. In general, the total heat released (THR) for PC– alumina nanocomposites did not change with an increase of alumina concentration, which can be explained by enhanced degradation processes catalyzed by alumina. Moreover, the time to ignition (TTI) generally decreases with the increase of alumina and/or alumina hydrate loading level. Totally different trends for the THR and the sample left were observed for PC–silica nanocomposites, as indicated in Table 7.7.12 The THR decreased and the amount of sample left increased with silica concentration increase, which indicated charforming mechanisms, as well as a radical trapping mechanism, by the silica surface. The restriction effects of silica as well as free-radical trapping mechanisms change the degradation pattern for PC. Free-radical degradation products, trapped by silica, could not leave the material, and the resulting free radicals are less reactive. PC–silica nanocomposites exhibited delayed TTI compared to PC. The restriction by silica of the PC chain-mobility effect combined with the free-radical trapping effect led to char formation.

7.3.2 PS–Inorganic Nanocomposites 7.3.2.1

Thermal Stability

Several factors affect the thermal stability of PS–alumina nanocomposites. The first is the stabilizing effect of alumina nanoparticles as restriction sites at the beginning of the thermal degradation process. Second, the degradation product free radicals are more reactive on the alumina oxide surface, which negatively affects the thermal stability of PS–alumina nanocomposites. Free-radical trapping on metal oxides increases the degradation rate, while chain-mobility restriction slows it down. It is a combination effect – PC and PS differ in that the catalysis effect does not occur in PS–alumina nanocomposites. Silica does not have a catalytic effect – it traps free radicals, which leads to a less reactive substance. The PS–silica degradation mechanism is associated with restriction of the PS chain mobility by silica nanoparticles. Therefore, the increase in silica concentration can result in higher degradation temperatures for PS–silica nanocomposites in comparison to that of PS. Burning tests are done to confirm the degradation. A small increase in the degradation temperatures for 10% and 50% weight loss was observed for PS–alpha-alumina (40–80 nm, spheres) nanocomposites and PS–alumina hydrate (100 nm/10 nm). Minimal increase in thermal stability was observed for PS–alumina nanocomposites with alumina concentration increase, because of the competing effects of accelerated degradation by free radicals and enhanced thermal stability caused by the PS chain-mobility restriction by

103

Impact of Nanoparticle Shape on the Flammability of Nanocomposites

alumina nanoparticles. Less enhancement occurred for PS–alumina (AL-15MAP and AL-wh-MAP) than for the PC counterparts at 10% weight loss. As shown in Tables 7.8 and 7.9,13 the thermal stability of PS–silica (16 nm, spherical) nanocomposites increased. This was attributed to trapping of freeradical species on the silica surface, which slowed the degradation process. Values at both 10 and 50% weight loss are higher for PS–silica nanocomposites than for PS–alumina nanocomposites. From Tables 7.8–7.10 it appears that the PS–inorganic nanocomposites with spherical nano-additives of larger size show better thermal stability compared to PS and a smaller size of nanoparticles. PS–10% silica 20 nm nanocomposites exhibited a 17 1C increase at 10% weight loss, while PS–10% silica 16 nm nanocomposites showed an 8 1C increase at 10% weight loss. From Table 7.10 it appears that the spherically shaped alumina of larger size (40–80 nm, alpha-form) exhibited better thermal stability than spherically shaped alumina hydrate (15 nm) or alumina hydrate nanorods, which can be attributed to the stronger interactions of PS chains with large-size spherical nanoparticles.

7.3.2.2

Flammability

With an alumina concentration increase, the burning rates measured by the horizontal burn (HB) test decreased, which means that PS–alumina burns Table 7.8

TGA results for PS–silica (16 nm) nanocomposites.13 10% weight loss

Sample PS (STYRON PS-Si-16-1 PS-Si-16-3 PS-Si-16-10 PS-Si-16-15

Table 7.9

s

685P)

50% weight loss

T (1C)

DT (1C)

T (1C)

DT (1C)

443 449 452 451 457

6 9 8 14

473 476 479 483 490

3 6 10 17

TGA results for PS–silica (20 nm) nanocomposites.14 10% weight loss

50% weight loss

Sample

T (1C)

DT (1C)

T (1C)

DT (1C)

PS (STYRONs 685D) PS-Si-20-PTCS-1 PS-Si-20-PTCS-3 PS-Si-20-PTCS-5 PS-Si-20-PTCS-10 PS-Si-20-PDMS-5 PS-Si-20-PDMS-10 PS-Si-20-PDMS-15

389 403 404 405 406 395 400 402

14 15 16 17 6 11 13

420 425 427 427 428 424 428 430

5 7 7 8 4 8 10

104

Table 7.10

Chapter 7

TGA results for polystyrene–alumina nanocomposites. 10% weight loss

Sample s

PS (STYRON 685DW) PS-Al-100-MAP-1 PS-Al-100-MAP-3 PS-Al-100-MAP-5 PS-Al-40-MAP-1 PS-Al-40-MAP-3 PS-Al-40-MAP-5 PS-Al-15-MAP-1 PS-Al-15-MAP-3 PS-Al-15-MAP-5 PS-Al-15-MT-1 PS-Al-15-MT-3 PS-Al-15-MT-5 PS-Al-wh-MAP-1 PS-Al-wh-MAP-3 PS-Al-wh-MAP-5 PS-Al-wh-MT-1 PS-Al-wh-MT-3 PS-Al-wh-MT-5

50% weight loss

T (1C)

DT (1C)

T (1C)

DT (1C)

385 389 386 389 389 392 392 384 391 381 387 387 389 385 392 393 387 390 368

4 1 4 4 7 7 –1 6 –4 2 2 3 1 7 8 2 5 –17

404 407 406 406 409 413 414 406 411 410 407 408 409 405 412 413 409 409 402

3 2 2 5 9 10 2 7 6 3 4 5 1 8 9 5 5 –2

longer than PS. In vertical burn (VB) tests, fire travels from bottom to top. Little to no time is available in which to form a protective barrier layer. Therefore, only the degradation-accelerating effect, caused by reactive freeradical products on the alumina surface, occur, which leads to more fuel supply for the burning process. Results of OI, VB and HB tests for PS–alumina nanocomposites are listed in Table 7.11. As discussed earlier, ignition and the continued burning of a material follow different mechanisms. The degradation mechanism only applies to the continued burning of the materials, such as in cone calorimetry. The OI test is a continued burning process that characterizes the ease of extinguishment for the sample. As can be seen in Table 7.11, the concentration or shape of alumina nano-additives had no significant effect on OI. The horizontal burning rate for PS–alumina nanocomposites generally decreased with increased concentration of alumina. PS–whisker alumina nanocomposites burned slower than PS (STYRONs 685 DW), as shown by the HB test (Table 7.11). PS–alumina nanocomposites burned faster in the VB test than PS. All PS–alumina samples dripped heavily in the HB test, which removed heat from the burning samples. In contrast, the burning rates for PS– silica (20 nm) nanocomposites were increased, as can be seen from Table 7.12.16 Materials became more flame retardant when both silica and brominated PS were used – the addition of silica lowered the content of brominated PS required to achieve the same level of flame retardancy, although, silica itself is not a flame retardant additive. The OI increased slightly for PS–silica (20 nm,

105

Impact of Nanoparticle Shape on the Flammability of Nanocomposites

Table 7.11

Burning tests results for polystyrene–alumina nanocomposites.

Sample s

PS (STYRON 685DW) PS-Al-100-MAP-1 PS-Al-100-MAP-3 PS-Al-100-MAP-5 PS-Al-100-MT-1 PS-Al-100-MT-3 PS-Al-100-MT-5 PS-Al-40-MAP-1 PS-Al-40-MAP-3 PS-Al-40-MAP-5 PS-Al-40-MT-1 PS-Al-40-MT-3 PS-Al-40-MT-5 PS-Al-15-MAP-1 PS-Al-15-MAP-3 PS-Al-15-MAP-5 PS-Al-15-MT-1 PS-Al-15-MT-3 PS-Al-15-MT-5 PS-Al-wh-MAP-1 PS-Al-wh-MAP-3 PS-Al-wh-MAP-5 PS-Al-wh-MT-1 PS-Al-wh-MT-3 PS-Al-wh-MT-5

Vertical burning test

OI

HB test, burning rate, (cm min
1)

T1 (s)

T2 (s)

Notes

16.9 16.9 17.2 17.5 16.9 17.2 17.5 17.2 17.5 17.2 17.1 17.1 16.9 17.5 17.2 17.2 17.5 17.2 17.5 17.5 17.5 17.2 17.5 17.2 17.2

6.1 5.8 5.2 4.4 5.5 5.1 5.1 5.8 5.3 5.4 6.2 6.3 5.4 4.4 5.4 5.2 5.4 5.1 4.7 6.1 4.4 4.5 4.7 4.7 3.9

76 69 53 66 59 54 55 58 62 51 61 63 61 60 55 47 53 43 21 41 55 61 59 49 58

– – – – – – – – – – – – – – – – – 7 11 – – – – – –

a a a a a a a a a a a a a a a a a a a a a a a a a

a

Ignites cotton and burns up to the clamp.

Table 7.12

OI and HB results for PS–silica (20 nm) nanocomposites.14,15

Sample s

PS (STYRON 685D) PS-Si-20-PTCS-1 PS-Si-20-PTCS-3 PS-Si-20-PTCS-5 PS-Si-20-PTCS-10 PS-Si-20-PTCS-15 PS-Si-20-PDMS-5 PS-Si-20-PDMS-10 PS-Si-20-PDMS-15 15%BrPS + 85%PS 30%BrPS + 70%PS 40%BrPS + 60%PS 15%BrPS + 75%PS + 10%Si 35%BrPS + 60%PS + 5%Si 15%BrPS + 84%PS + 1%Si a

OI

HB test, burning rate, (cm min
1)

17.3 17.0 17.3 17.5 18.1 18.4 17.5 17.8 18.1 18.9 21.4 24.0 20.5 25.0 19.6

5.4 5.7 5.8 6.1 7.3 7.1 6.6 8.2 9.3 4.1 b AEB ¼ 3.5 cm b AEB ¼ 0.5 cm b AEB ¼ 6.5 cm b AEB ¼ 3 cm 4.3

Ignites cotton and burns up to the clamp. EB ¼ average extent of burning.

b

Vertical burning test T1 (s)

T2 (s)

104 92 83 68 57 – 49 53 62 8 9 1 63 2 5

– – – – – – – – – 55 5 1 – 3 50

Notes

–

a

V2 V2

a

V2

a

106

Table 7.13

Chapter 7 s

Cone calorimetry results for PS (STYRON 685P)–silica (16 nm, spherical) nanocomposites.13

Sample PHRR (kW m
2) Time of PHRR (s) THR (MJ m
2 g
1) Average mass loss rate (g m
2 s–1) Sample left (%)

PS(STYRONs 685P)

PS-Si-16-3

PS-Si-16-10

1361 69.0 3.42 39.6

1178 (–13.4%) 49.2 3.27 (–4.2%) 34.8

1048 (–22.1%) 42 2.97 (–13.1%) 30.9

0.0

3.2

12.6

spherical) nanocomposites compared to PS (STYRONs 685D). The horizontal burning rates increased with silica concentration increase for PS–silica (20 nm, spherical) nanocomposites. Of note is that PS with 3% and 5% of MT-treated whiskers in Table 7.11 are not V-2 materials, since they all burn to the clamp. Severe dripping behaviour is the cause of flame extinction in the first flame application. In cone calorimetry, each sample is subjected to a slow burning process. Therefore, there is enough time for free-radical product trapping by silica to occur, which slows the degradation process. The free-radical trapping effect, caused by silica, leads to lower PHRR and THR, as can be seen in Table 7.13.13 The amounts of sample left for PS–silica nanocomposites correspond to the concentration of silica present in the samples, which indicates that char formation did not occur.

7.3.2.3

Mechanical Properties

Results of mechanical tests indicate the strength of interfacial interactions for polymer–inorganic nanocomposites. Polymers with different stereochemistry are expected to interact with different shapes of nanoparticles with different efficacy. Here, the following simplified model is proposed for the reinforcement mechanisms in polymer–inorganic nanocomposites. The aromatic ring of polymers has a flat architecture; therefore, the area of interaction for this type of polymer with the one-dimensional additives (nanorods) is larger than with zerodimensional ones (nano-spheres). Stronger interfacial interactions are expected between aromatic polymers and one-dimensional additives. As for polymers with a flexible structure, they can entangle around spherical nanoparticles, and strong interfacial interactions will occur. PMMA–silica nanocomposites exhibited a 125% increase in tensile strength at 5% loading (Table 7.12), while PS–silica and PC–silica nanocomposites showed no improvement in tensile strength and elongation at break. In contrast, PC–alumina nanocomposites from alumina nanorods and alumina whiskers showed substantial improvement in the mechanical performance compared to PC (54% increase at 2% loading for alumina hydrate 100/10 nm) Table 7.14.

107

Impact of Nanoparticle Shape on the Flammability of Nanocomposites

Table 7.14

Mechanical properties of polymer–silica and polymer–alumina nanocomposites from extrusion.13,16

Sample

Tensile strength (MPa)

Modulus (MPa)

Elongation at break (%)

PMMA PMMA-20nm silica-5% PS(685) PS-20nm silica-5% PC (Lexan) PC-16nm silica-!% PC-16nm silica-3% PC-16nm silica-5% PC (Dow) PC-40nm alumina-1% PC-40nm alumina-2% PC-40nm alumina-5% PC-100nm alumina hydrate-1% PC-100nm alumina hydrate-2% PC-100nm alumina hydrate-5% PC (OQ 3620 110) PC-alumina whisker-1% PC-alumina whisker-3% PC-alumina whisker- 5%

16.5 37.2 33.1 31.7 56.5 57.2 58.6 57.2 59.9 79.7 76.6 50.9 81.2 92.2 61.2 65.1 84.1 82.8 78.9

178 258 557 869 304 291 323 392 194 315 322 314 350 382 418 319 393 409 433

8 17 10 10 67 62 69 36 117 78 102 25 49 30 21 60 78 44 29

Given the differences in interfacial interaction between different polymer matrixes, nanoparticle chemistry and shapes result in different flammability properties for nanocomposites. The shapes of nanofillers may greatly affect the degree of interfacial interaction between polymer matrix and nanoparticles, which in turn significantly impacts the thermal stability and flammability of nanocomposites. The largest tensile strength increase for PC is with 2% 100 nm alumina. That also has the largest reduction in PHRR at –46%.

7.4 Conclusion It was found that the chemical nature of nanoparticles and polymers played an important role in determining the mechanisms of degradation. Metal oxide nanoparticles (alumina) have the ability to catalyze the degradation of carbonyl-containing polymers as well as to trap free radicals, which accelerates the degradation process. It was found that alumina moved to the surface and formed a protective or ‘‘barrier’’ layer, which enhanced thermal stability for polymer–alumina nanocomposites. This layer slows the processes of mass and energy transfer to and from the material. Polymer–silica nanocomposites showed higher degradation temperatures, but the degradation mechanism is different for polymer–alumina nanocomposites. Since silica is not a metal oxide, it does not have a catalytic effect, and the free-radical trapping on the

108

Chapter 7

surface of silica leads to less-reactive free radicals. Moreover, silica forms restriction sites for polymer chain mobility, which slows the degradation processes for polymer–silica nanocomposites. It was found that the catalytic and free-radical trapping effects of alumina accelerated the burning process, while the formation of an alumina ‘‘barrier’’ layer enhances the flame retardancy for polymer–alumina nanocomposites. In contrast, free-radical trapping and restriction effects of silica improve flame retardancy of polymer–silica nanocomposites. Silica enhances char formation for char-forming polymers (for instance, PC), which enhances the flame retardancy of char-forming polymer–silica nanocomposites. It was shown that PC and PS matrixes have stronger interactions with a one-dimensional additive (nanorod) compared to a zero-dimensional additive (nanosphere). Thermal degradation and flammability for polymer–inorganic nanocomposites from different shapes of nano-additive exhibited quite different behaviour, as demonstrated by cone calorimetry, OI, HB and VB test results.

References 1. P. Rittigstein and J. M. Torkelson, J. Polymer Sci. Part B: Polymer Physics, 2006, 44, 2935–2943. 2. V. K. Nguyen, J. W. Lee and Y. Yoo, Sensors and Actuators B: Chemical, 2007, 120, 529–537. 3. S. Chatterjee, A. Goyal and S. I. Shah, Mater. Lett., 2006, 60, 3541–3543. 4. C. E. Powell and G. W. Beall, Current Opinion in Solid State and Materials Science, 2006, 10, 73–80. 5. F. Yang and G. L. Nelson, The 11th International Conference, ADDITIVES, 2002. 6. J. Zhu, F. M. Uhl and C. A. Wilkie, Polymer Preprints, 2001, 42, 392. 7. M. Moniruzaman and K. J. Winey, Macromolecules, 2006, 39, 5194–5205. 8. Y. W. Mai and Z. Z. Yu, ‘‘Polymer Nanocomposites’’, 2006, Woodhead Publishing Limited, Cambridge, 57–100. 9. S. C. Liufu, H. N. Xiao and Y. P. Li, Polymer Degradation and Stability, 2005, 87, 103–110. 10. A. Laachachi, M. Cochez, E. Leroy, P. Gaudon, M. Ferriol and J. M. Lopez Cuesta, Polymers for Advanced Technologies, 2006, 327–334. 11. I. C. McNeil and H. M. Mussarat, Polymer Degradation and Stability, 1995, 48, 189–195. 12. F. Yang, R. Yngard, A. Hernberg and G. L. Nelson, Proceedings of BCC 16th Annual Flame Conference, 2005, 244–255. 13. F. Yang and G. L. Nelson, Polymers for Advanced Technologies, 2006, 17, 320–326. 14. R. A. Yngard, ‘‘Preparation and characterization of flame retardant polystyrene/silica nanocomposites’’, A thesis submitted to Florida Institute

Impact of Nanoparticle Shape on the Flammability of Nanocomposites

109

of Technology in partial fulfillment of the requirements for the degree of master of science in chemistry, 2002, Melbourne, Florida. 15. F. Yang, R. Yngard and G. L. Nelson, J. of Fire Sci., 2005, 23, 209–226. 16. I. R. Bogdanova, ‘‘Reinforcement and degradation mechanisms in polymer/inorganic nanocomposites’’, 2007, Dissertation submitted to Florida Institute of Technology in partial fulfillment of the requirements for the degree of Doctor of Philosophy in Chemistry, pp. 61, 92.

CHAPTER 8

Thermal and Combustion Behaviour of Polymer–Carbon Nanofibre Composites D. TABUANI, S. PAGLIARI, W. GIANELLI AND G. CAMINO Centro di Cultura per l’Ingegneria delle Materie Plastiche, viale T. Michel 15, Alessandria, Italy

8.1 Introduction In recent years the use of graphitic nanoparticles, such as carbon nanofibres (CNFs) and nanotubes, has attracted much interest for the development of highperformance polymer composites. In particular, CNFs are an attractive alternative to nanotubes, because of their lower costs while maintaining properties similar to those of their thinner analogues. Indeed, the use of CNFs as polymer fillers is quite recent,1,2 but has developed quite fast. In particular, studies have been carried out with polypropylene (PP) in which nanofibres were found to improve thermal stability,1,3 crystallization behaviour,1,4 mechanical properties,1,4,5,6 melt viscosity,2 and electrical2 and thermal conductivity.6 Research has been devoted also to other polymer classes, such as poly(ether ether ketone)7,8, polycarbonate,6 poly(methyl methacrylate),9 epoxy10 and polystyrene (PS).11,12 With particular reference to fire retardancy, nanoadditives have recently attracted attention for their potential as fire retardants of low environmental impact, as alternatives for the more commonly used but more hazardous aromatic brominated compounds. This aspect has been developed for layered silicates,13 and also carbon nanoparticles were found to have a good potential Fire Retardancy of Polymers: New Strategies and Mechanisms Edited by T Richard Hull and Baljinder K Kandola r Royal Society of Chemistry 2009 Published by the Royal Society of Chemistry, www.rsc.org

110

Thermal and Combustion Behaviour of Polymer–Carbon Nanofibre Composites

111

in the field. In particular, carbon nanotubes were found to improve significantly the fire retardant behaviour of ethylene vinyl acetate (EVA),14 PP3,15,16 and polyamide 6 (PA6)17 alone and in combination with layered silicates18. More recently this aspect was verified also for CNFs in PP, using the limiting oxygen index (LOI) as the flammability test.3 This chapter describes our investigation of the influence of the addition of CNFs on the properties of two polymer matrices, one semicrystalline and the other completely amorphous (PP and PS, respectively). Furthermore, two grades of PP are taken into account, with different melt-flow indexes (MFI), to investigate the influence of PP melt viscosity on the properties of the final composite. In the case of PS, preparation of the composites by melt blending is compared with a solution method. Two different kinds of nanofibres with different aspect ratios are considered, to evaluate the influence of this parameter on the fibre dispersability and final material properties. The morphology, thermal and combustion behaviour of the composites are evaluated. This chapter represents the first attempt to evaluate the influence of polymer characteristics, preparation conditions and nanofibre type on the final material properties.

8.2 Materials and Methods Two PP grades were used, both purchased from Basell, Italy: Moplen HP400R (MFI ¼ 26 g 10 min
1) and Moplen HP501L (MFI ¼ 6 g 10 min
1). PS (EDISTIR N 2380) was supplied by Polimeri Europa (Italy). Nanofibres, PR-24 HHT, were purchased from Pyrograf Products (USA) in two commercial grades: high density (HD) and low density (LD) standards. The HD and LD designations do not refer to the density of the individual fibre but to the bulk density of the CNFs. The bulk or aggregate density of the nanofibre is controlled by altering the intensity and duration of the process that reduces nanofibre aggregation, which also reduces the length of the individual fibres, with HD fibre having shorter lengths (50–100 mm) than the low density fibre (100–200 mm). The fibre diameter, conversely, is not influenced by the process, ranging in any case from 100 to 200 nm. Two techniques were used to prepare the composites: polymer–fibre melt blending with PP and PS, and dispersion of fibres in polymer solution in the case of PS.

8.2.1 Melt Blending Nanofibres were mechanically dispersed in the molten polymer by the shear action of a Brabender internal mixer, model PLE 67152. The mixing parameters for each matrix are reported in Table 8.1.

8.2.2 From Solution PS (4 g) was dissolved with a magnetic stirrer in 15 cm3 of tetrahydrofuran (THF, Aldrich) at room temperature and the solution obtained was mixed with

112

Table 8.1

Chapter 8

Melt processing conditions.

Matrix

Process temperature (1C)

Screw rotation speed (rpm)

Residence time (minutes)

PPHP400R

180

60

10

PPHP501L

180

60

10

PS

180

30

5

Loading (wt%) LDCNF (0.5; 1; 3) HDCNF (0.5; 1; 3) LDCNF (0.5; 1; 3) HDCNF (0.5; 1; 3) LDCNF (3; 6; 10) HDCNF (3; 6; 10)

5 cm3 of nanofibre suspension in THF (the concentration of the suspension varied depending on the filler loading), using a magnetic stirrer to break-up any nanofibre macro-agglomerates. The successive dispersion of nanofibre suspension in the polymer solution was achieved with an ultrasound bath model Uniset AC14 (power, 140 W; frequency, 22 kHz). The solvent was eliminated under vacuum in a Rotavapor, followed by vacuum oven drying (T ¼ 80 1C, P ¼ 3 mbar, t ¼ 180 minutes).

8.2.3 Characterization X-Ray (WAXRD) diffraction patterns were obtained on a ARL XTRA48 diffractometer using Cu Ka radiation (l ¼ 1.54062 A˚) on compression moulded 1 mm thick specimens. The crystallization behaviour of the PP samples was measured using a TA Q1000 instrument in hermetically sealed aluminium pans under nitrogen flow (50 cm3 min
1). Three successive runs (heating–cooling–heating) were performed at 10 1C min
1 from 30 to 250 1C, with ca. 5 mg samples. The first heating run erases the thermal history of the samples and an annealing process at 250 1C for 10 minutes followed the first heating run. A detailed differential scanning calorimetry (DSC) analysis was performed on PP nanocomposites to assess the degree of crystallinity and crystallization rate, measuring crystallization temperature (Tc), melting enthalpy (DHm) and degree of crystallization (Xc). Crystallization temperature was measured within a  0.5 1C tolerance. The degree of crystallization was calculated from the peak enthalpy area normalized to the actual polymer weight fraction according to:

Xc ¼

DHm  100 DHm0  Wpolymer

in which DHm is the theoretical value of enthalpy for 100% crystalline polymer (190 J g
1)19 and Wpolymer is the weight of the polymer fraction. Thermogravimetric analysis (TGA) was carried out both in inert (nitrogen) and in oxidizing (air) atmospheres on a TA Q500 instrument on ca. 10 mg samples, in platinum pans, with gas fluxes of 60 cm3 min
1 for sample gas

Thermal and Combustion Behaviour of Polymer–Carbon Nanofibre Composites 3

113


1

(nitrogen or air), and 40 cm min for balance-protection gas (nitrogen) at a fixed rate (in this case 10 1C min
1) from 50 to 800 1C. From the thermograms two relevant parameters were taken to enable discussion: T5%, as the temperature at 5% of weight loss, and Tmax, as the temperature at which the maximum weight-loss rate occurred (i.e. derivative TGA, peak temperature). T5% and Tmax were measured within a  3 1C tolerance. Combustion tests were performed on a Fire Testing Technology Cone Calorimeter, with 50  50  3 mm specimens, prepared by compression moulding on a hot-plate laboratory press at 180 1C. Tests were performed at 50 kW m
2 external heat flux, to evaluate the fire properties of the composites in conditions comparable to a developed fire scenario. Specimens were wrapped in aluminium foil (except for the upper surface) and placed on a ceramic backing board at a distance of 25 mm from the cone base. The average values of three successive experiments are discussed here. Scanning electron microscopy (SEM) images were obtained by means of a LEO 1450 VP instrument on cryogenically fractured surfaces. For comparison, characterizations were also performed on pure PP and PS treated in the same way as the composites.

8.3 Results and Discussion 8.3.1 Morphology 8.3.1.1

SEM

SEM images show fairly good dispersion of nanofibres in the polymer matrices. In particular, as far as PP with HDCNF samples are concerned, nanofibre-rich areas can be identified at low magnification [Figure 8.1(a)], but single nanofibres well-adhered to the polymer matrix are also visible [Figure 8.1(b)]. Similar results are obtained with LDCNF and with PPHP501H. The dispersion of the fibres in PS is even better; indeed, these appear to be uniformly distributed in the matrix [Figure 8.2(a)], and also present a high degree of compatibility with the polymer itself [Figure 8.2(b)].

(a)

Figure 8.1

(b)

SEM micrographs of HP400R–3wt%HDCNF composites.

114

Chapter 8

(a)

Figure 8.2

(b)

SEM micrographs of PS–3wt%HDCNF composites from melt blending.

(a)

Figure 8.3

(b)

SEM micrographs of PS–3wt%HDCNF composites from solution.

The situation appears to be dramatically different when nanofibres are dispersed in PS via the solution method. The mixing technique was apparently not efficient in the disaggregation of the nanofibre bundles, which are observed in the polymer matrix even at low magnification [Figure 8.3(a)]. When observing the area in detail we could recognize the aggregates as composed only of nanofibre – the polymer, in this case, is not interacting at all with the filler.

8.3.1.2

XRD

In XRD, diffractograms for the PPHP400R–HDCNF composite, Figure 8.4, the most intense reflection peak of pristine CNF is visible at 2y ¼ 25.5. The PP matrix in the composite crystallizes with the same phase as that of the pristine polymer, as shown by comparison with XRD of pure PP (Figure 8.4). However, a strong (040) preferential orientation is observed, owing to the presence of the nanofibres, even at the lowest loading. Indeed, the significant increase of the intensity ratio between the second (040 reflection, 2y ¼ 17) and the first (110 reflection, 2y ¼ 14.3) XRD peak results from a preferential growth

Thermal and Combustion Behaviour of Polymer–Carbon Nanofibre Composites

115

2θ = 25.5

Intensity (a.u)

PPHP400R+3%HDCNF

PPHP400R+1%HDCNF

PPHP400R+0,5%HDCNF (110) (040)

6

Figure 8.4

(130) (111) (131) (041)

PPHP400R

8 10 12 14 16 18 20 22 24 26 28 30 32 34 36 38 40 42 44 2θ

XRD diffraction pattern of HP400R–HDCNF.

of PP crystallites well-ordered along the b direction. This phenomenon was reported in the literature for other nanofillers.20–22 The hypothesis is that this preferential orientation is generated by an intimate interaction of the polymer chains with the dispersed nanofiller, so that the filler can influence the growth direction of the crystalline domains. We could not find significant differences for this aspect when considering the two PP grades and CNF types. In Figure 8.5, XRD spectra of blank PS and LDCNF-containing composites prepared by melt blending are compared. A wide peak from 2y ¼ 15 to 251 belongs to amorphous PS and the most intense reflection peak of CNF is visible in the spectra of the composites at 2y ¼ 25.5.

8.3.2 Thermal Behaviour 8.3.2.1

DSC

Temperatures of crystallization and degree of crystallinity measured during the second heating of nanocomposite materials are reported in Table 8.2. Nanofillers were observed to act as nucleation sites, as evidenced by a significant decrease in the time required by nanocomposites to crystallize during the cooling cycle. As a consequence, the nanocomposite crystallization peak endotherm temperature was higher than the value observed for the pure

116

Intensity (a.u)

Chapter 8

PS+10%LDCNF PS+6%LDCNF PS+3%LDCNF PS 0 4 8 12 16 20 24 28 32 36 40 44 48 52 56 60 64 68 2θ

Figure 8.5

Table 8.2

XRD diffraction pattern of PS–LDCNF composites.

Crystallization temperatures (Tc) and degree of crystallinity (Xc) for CNF–PP composites.

HP400R

Tc (increment) (1C)

Xc (variation) (%)

PPHP501L

Tc (increment) (1C)

Xc (variation) (%)

– +0.5%LDCNF +1%LDCNF +3%LDCNF +0.5%HDCNF +1%HDCNF +3%HDCNF

109 123 124 128 121 122 125

48 49 54 52 49 51 50

– +0,5%LDCNF +1%LDCNF +3%LDCNF + 0,5%HDCNF + 1%HDCNF + 3%HDCNF

111 123 124 128 121 122 124

48 54 51 50 50 51 50

(+14) (+15) (+19) (+12) (+13) (+16)

(+1%) (+5%) (+4%) (+1%) (+3%) (+ 2%)

(+12) (+13) (+17) (+10) (+11) (+13)

(+6%) (+3%) (+2%) (+2%) (+3%) (+2%)

polymer, as shown in Figure 8.6 for PPHP400R. The same behaviour was observed for PPHP501L. Data on the crystallization behaviour of both PP with HD and LD nanofibres are given in Table 8.2. No significant differences were encountered when comparing LD- and HDCNF and nor were any notable variations of Tc, found when increasing the content from 1 to 3%, which thus indicates that saturation of the crystallization nucleation sites was reached at 1% loading. This phenomenon has already been observed by Lozano et al.1 and was attributed by them to the increased number of nucleating sites, thus indicating that CNFs act as heterogeneous nucleating agents for PP. This phenomenon is in accordance with that observed from XRD spectra, which suggests a close interaction between the nanodispersed phase and the polymer chains. As evident from Table 8.2, PP crystallinity is slightly increased when there are nanofibres in the matrix; other authors have already reported such behaviour for

Thermal and Combustion Behaviour of Polymer–Carbon Nanofibre Composites

117

PP400R PP400R+0,5%HDCNF PP400R+1%HDCNF PP400R+3%HDCNF

18 16 14

Heat flow (a.u.)

12 10 8 6 4 2 0 80

100

120

140

160

Temperature (°C)

Figure 8.6

Specific heat flow (cooling step) recorded for HP400R–HDCNF composites.

similar systems1 in which vapour-grown CNF–PP composites were prepared, and the influence of degree of purification and of functionalization of the nanofibres on PP thermal and mechanical properties was determined. In the work by Lozano et al., at a concentration of 5% CNF, an increase in PP crystallinity of about 4% was observed and attributed to an enhanced mobility of the macromolecular chains, which enables a better alignment of the crystal lattice.

8.3.2.2

TGA – Polypropylene Based Materials

During heating under nitrogen, PP thermally degrades in a single step to volatile products above 350 1C (T5% 392 1C, Tmax 450 1C, dotted line in Figure 8.7) through a radical chain process propagated by carbon-centred radicals that arise from carbon–carbon bond scission.23 The thermal degradation of PP that contains CNFs takes place in a single step, as in the case of pristine PP, and shows the same temperature for a 50% weight loss. In air, above 200 1C the radical chain thermal volatilization is initiated by H abstraction from PP by oxygen.23 Volatilization begins in TGA below 250 1C (T5%), with Tmax at 324 1C (solid line in Figure 8.7; see Table 8.3), and is completed before the temperature of the initiation of the pure thermal degradation process (370–390 1C) is reached. The presence of CNF modifies the thermo-oxidative behaviour of the matrices in all the prepared nanocomposites, increasing the temperature at which 50% volatilization occurs, as shown in Figure 8.7 and Table 8.3.

118

Chapter 8 277(95%)

100

253(95%)

80 Weight [%]

PPHP400R PPHP400R+0,5% LDCNF PPHP400R+1% LDCNF PPHP400R+3% LDCNF PPHP400R nitrogen

392°C (95%)

60 40 20

Deriv. Weight [%/˚C]

0 377°C 2

450°C

324°C

1

0 100

Figure 8.7

Table 8.3

200

300

400 500 Temperature [°C]

600

700

800

TGA and differential thermogravimetric analysis (DTG) curves in air of HP400R–LDCNF composites.

Degradation temperatures for CNF–PP composites in air.

HP400R

T5% (1C)

Tmax (1C)

HP501L

T5% (1C)

Tmax (1C)

– 0.5%LDCNF 1%LDCNF 3%LDCNF 0.5%HDCNF 1%HDCNF 3%HDCNF

253 255 250 277 250 253 261

324 348 352 377 334 332 361

– 0.5%LDCNF 1%LDCNF 3%LDCNF 0.5%HDCNF 1%HDCNF 3%HDCNF

255 258 255 277 257 266 272

328 365 359 382 320 359 371

We attribute the protective action played by the nanofibres to a barrier effect against oxygen, provided by the nanofibres which organize themselves in a physical network during polymer ablation, and to an interaction of the fillers with oxygen. Moreover, this physical network could provide a mass-transport barrier towards polymer ablation. Both barrier effects lead to reduction of the observed mass-loss rate. Indeed, an interesting feature is, for both PP grades investigated and for both nanofibres, the shape of the derivative curves, i.e. the shape of the mass-loss rate. It can be seen that the pure polymer derivative curve differs significantly from those of the filled samples. At low temperatures, up to 320 1C,

Thermal and Combustion Behaviour of Polymer–Carbon Nanofibre Composites

119

the mass-loss rate of the PP in the composites is significantly lower than that of the pristine polymer (lower slope of the curve, especially for the 3% loaded samples). At above 320 1C an increase of the mass-loss rate occurs up to the maximum value (identified for the various composites by Tmax). Above Tmax an abrupt increase of the mass-loss rate is experienced by the filled samples, the overall result being a completely asymmetrical peak. We can recognize in this behaviour the efficiency of the network described above, which is high at low temperatures and decreases upon heating at higher temperatures. Above Tmax the physical rupture of this network occurs: the remaining polymer experiences a combination of high temperature and oxygen attack that makes it degrade very rapidly (the derivative curve is almost vertical above Tmax).

8.3.2.3

TGA – Polystyrene Based Materials

PS is known to degrade through a radical mechanism in nitrogen, forming styrene as a major product via b-scission of chain-end radicals; dimers, trimers and larger oligomers are also produced via intramolecular hydrogen transfer reactions (backbiting) with the yield decreasing with oligomer size.24 As far as PS composites prepared from the solution method are concerned, a low temperature peak is visible on the DTG curves (ca. 150 1C, both in nitrogen and in air), which arises from residual solvent from the preparation method. This signal is more intense for the composites than for the pure material, probably owing to the lower tendency of the solvent to evaporate from a more viscous medium or to the barrier effect towards evaporation played by the nanofibres. The nanofillers did not have an effect on the single-step weight-loss thermal degradation process of PS. Whereas under air (Figure 8.8), in which oxygen initiation decreases the temperature of weight loss (as in PP) although with a lower effect (Tmax N2 415 1C, Tmax air 374 1C), a delay in the degradation temperatures is observed with 3% of filler and further addition of nanofibre does not produce any relevant changes. The maximum degradation-rate temperature of PS becomes closer to that of the pure polymer in nitrogen, which indicates an effective protection role played by the fibres towards the degradative action of oxygen, which abstracts hydrogen and reacts with polystyryl radicals that accelerate the overall process of weight loss.24 As in the case of PP, also in PS, above a critical temperature (ca. 475 1C for PS), the degradation rate rapidly increases, probably because of the breakdown of the nanofibre protecting network. For composites prepared by melt blending, a 5 1C increase in Tmax is observed for the degradation in nitrogen. In air, in contrast, a significant increase in the degradation temperature is observed, independent of the amount of filler; such an increase is more noticeable than that observed in the samples from solution, probably because of a better miscibility given by the preparation method. Indeed, especially in the case of HDCNF, the maximum weight-loss rate temperature of the matrix in nitrogen is achieved, thus implying a very efficient protection action towards oxygen (Figure 8.9).

120

Chapter 8

Weight [%]

80

PS PS+3% LDCNF PS+6% LDCNF PS+10% LDCNF PS nitrogen

309°C(95%)

100

370°C (95%) 151°C(95%)

60 40 20 0 402°C

Deriv. Weight [%/˚C]

3,0

416°C

2,5 2,0 374°C

1,5 1,0 0,5 0,0 100

Figure 8.8

8.3.2.4

200

300

400 500 Temperature [°C]

600

700

800

TGA and DTG curves in air of PS–LDCNF prepared by the solution method.

Cone Calorimeter Test

The reaction to fire of polymer matrices and nanocomposite materials was evaluated by the cone calorimeter test and, in particular, heat release rate (HRR) patterns were recorded. The time to ignition (TTI), the maximum value of HRR plotted versus time [peak HRR (PHRR)] and the total amount of heat released during the combustion process [total heat release (THR)] are the parameters discussed here. In the case of nanocomposites based on PP it is evident, from Table 8.4 and Figure 8.10, that TTI decreases with both LD and HDCNF through an increased radiant-heat absorbance compared to that of pure PP because of the black colour of the nanocomposites. For PP composites the total heat released during combustion is maintained in the same range. To reduce appreciably the PHRR of PP (from ca. 1330 to 890 kW m
2) 3% of LDCNF has to be added to PP. Lower concentrations did not modify PHRR. HDCNFs are less efficient because, even at 3% loading, they did not affect PHRR. Similar results are obtained with PPHP501L. We attribute the effect induced by LDCNF to the assembly of a protective CNF network which lowers the combustion rate, in analogy with that described for thermo-oxidation, which does act as a barrier towards the release of combustible volatile products.

Thermal and Combustion Behaviour of Polymer–Carbon Nanofibre Composites 324°C (95%) 377°C (95%) 306°C (95%)

100

Weight [%]

80

121

PS PS+3% LDCNF PS+6% LDCNF PS+10% LDCNF PS nitrogen

60 40 20 0

Deriv. Weight [%/˚C]

3

414°C

2 375°C

411°C

1

0 100

200

300

500 600 400 Temperature [°C]

700

800

Figure 8.9

TGA and DTG curves in air of PS–LDCNF prepared by melt blending.

Table 8.4

Cone calorimeter results for nanocomposites based on PPHP400R.

HP400R

TTI (s)

PHRR (kWm–2)

THR (MJ m–2)

– +0.5%LDCNF +1%LDCNF +3%LDCNF +0.5%HDCNF +1%HDCNF +3%HDCNF

37  5 20  0 24  2 23  3 23  2 21  2 23  2

1331  40 1356  20 1266  35 891  15 1353  65 1400  20 1237  30

93  7 99  0 99  1 95  2 97  2 99  2 98  2

In all the cases described, the weight of residue obtained after the combustion process was equivalent to the weight of the employed CNF, which rules out any extensive charring effect of CNFs. For the PS system, the cone calorimeter analysis was performed on the meltblended samples at 3 and 10% loadings to evaluate the influence on the polymer fire behaviour of the lowest and highest nanofibre percentages in the composites that showed the best CNF dispersion. As shown in Figure 8.11, the two kinds of nanofibres behave differently. As far as HDCNFs are concerned, an appreciable decrease in the PHRR (about 50%) was obtained by adding 10% of filler, whereas no effect was detected at 3% loading, within the experimental error (Table 8.5).

122

Chapter 8 PPHP501L PPHP501L+0,5% LDCNF PPHP501L+1% LDCNF PPHP501L+3% LDCNF

1400

600

2

(a)

800

HRR (KW/m )

1200

1000

2

HRR (kw/m )

1200

1000

(b)

800 600

400

400

200

200

0

0 −20 0 20 40 60 80 100 120 140 160180 200 220 240 time (sec)

0 20 40 60 80 100 120 140 160 180 200 220 240 time (sec)

Figure 8.10

1600 1400 1200 1000 800 600 400 200 0

HRR vs. time: (a) HP501L–LDCNF and (b) HP501L–HDCNF composites.

(a)

0

Figure 8.11

Table 8.5

PS PS+3% LDCNF PS+10% LDCNF HRR (kW/m2)

HRR (kW/m2)

PPHP501L PPHP501L+0,5% HDCNF PPHP501L+1%HDCNF PPHP501L+3%HDCNF

1400

50

100 150 200 250 time (sec)

PS PS+3% HDCNF PS+10% HDCNF

(b)

1600 1400 1200 1000 800 600 400 200 0 0

50 100 150 200 250 300 time (sec)

HRR vs. time: (a) PS–LDCNF and (b) PS–HDCNF composites.

Cone calorimeter results for nanocomposites based on PS.

PS

PHRR (kW m 2)

TTI (s)

THR (MJ m
2)

– +3%LDCNF +10%LDCNF +3%HDCNF +10%HDCNF

1536  50 977  40 858  35 1337  40 758  30

43  1 31  3 39  2 25  4 32  2

91  1 88  1 88  4 91  1 88  3

Conversely, with LDCNF a significant decrease of the PHRR value is obtained by the addition of the lowest amount of filler (3%) and the increase of such loading (to 10%) does not bring significant further improvements on the fire behaviour of the material. Indeed, by the addition of 3% LDCNF, a decrease by 36% of the PHRR value was observed and, by increasing the LDCNF content by a factor of three (to 10%), the PHRR decreases with

Thermal and Combustion Behaviour of Polymer–Carbon Nanofibre Composites

123

respect to that of pristine polymer by about 44%, that is 20–25% more than with the 3% loading.

8.4 Conclusions No differences were encountered in the miscibility of LD or HD nanofibres with PP, either of high or low molecular weight. In all cases, CNFs were found to be fairly well dispersed in the polymer and to behave as nucleating agents, increasing the polymer crystallization temperature and influencing, in the same way, a preferential orientation of PP chains along the b direction during crystallization. Significant improvements were found in PP thermo-oxidative resistance, to a greater extent in the case of LDCNF. Similarly, LDCNF shows a greater efficiency than HDCNF in reducing the PHRR of PP burning in a cone calorimeter. In PS-based materials, a noticeable effect on the properties of the composites was played by the preparation technique, mainly because of the different degree of dispersion achieved with the nanofillers; in fact, the solution method leads to an extremely poor dispersion of the CNF in the polymer. Nevertheless, also in this case, some improvements in PS thermo-oxidative behaviour were observed, yet the most interesting results were obtained by adding the filler via the meltblending method. In this case, indeed, the addition of HDCNF in particular shifted PS degradation temperatures towards those observed in nitrogen, thus proving the efficiency of the nanofibres in protecting the polymer from the degradative action of oxygen. As far as combustion tests are concerned, a significant influence of the type of filler was found on the fire behaviour of PS; LDCNFs were found to behave efficiently as fire retardants even at the lowest loading investigated (3%), while a significant decrease in the PHRR was found only by adding 10% of HDCNFs. It is clear, therefore, that specific interactions are occurring between the two polymers investigated and LDCNF, making this the most suitable filler to tune the properties of the polymer. We can hypothesise that, because LDCNF has a higher aspect ratio than HDCNF, a greater interface with the polymer is created, and it is this that is responsible for the better performance. Recently, the hypothesis of different interaction extents was confirmed for the PP matrix (in particular, for the HP400R grade) by means of rheological measurements.25

References 1. K. Lozano and E.V. Barrera, J. Appl. Polym. Sci., 2001, 79, 125–133. 2. K. Lozano, J. Bonilla-Rios and E.V. Barrera, J. Appl. Polym. Sci., 2001, 80, 1162–1172. 3. A. Chatterjee and B.L. Deoupra, J Appl Polym Sci, 2006, 100, 3574–3578. 4. X. Tong, Y. Chen and H. Cheng, J. Mater. Sci. Technol., 2005, 21, 686. 5. S. Kumar, H. Doshi, M. Srinivasarao, J.O. Park and D.A. Schiraldi, Polymer, 2002, 43, 1701–1703.

124

Chapter 8

6. E. Hammel, X. Tang, M. Trampert, T. Schmitt, K. Mauthner, A. Eder and P. Poetschke, Carbon, 2004, 42, 1153. 7. J. Sandler, P. Werner, M.S.P. Shaffer, V. Demchuk, V. Altstadt and A.H. Windle, Composites: Part A, 2002, 1033. 8. J. Sandler, A.H. Windle, P. Werner, V. Altstadt, M.V. Es and M.S.P. Shaffer, J. Mater. Sci., 2003, 38, 2135. 9. J. Zeng, B. Saltysiak, W.S. Johnson, D.A. Schiraldi and S. Kumar, Composites Part B, 2004, 35, 245. 10. Y.K. Choi, K. Sugimoto, S.M. Song, Y. Gotoh, Y. Ohkoshi and M. Endo, Carbon, 2005, 43, 2199. 11. B. Zhang, R. Fu, M. Zhang, X. Dong, L. Wang and C. Pittman, U. Mater. Res. Bull., 2006, 41, 553. 12. Y. Xu, B. Higgins and W.J. Brittain, Polymer, 2005, 46, 799–810. 13. A.B. Morgan, Polym. Adv. Technol., 2006, 17, 206–217. 14. G. Beyer, Fire Mat., 2002, 26, 291. 15. T. Kashiwagi, E. Grulke, J. Hilding, R. Harris, W. Awad and J. Douglas, Macromol Rapid Commun, 2002, 23, 761–765. 16. T. Kashiwagi, E. Grulke, J. Hilding, K. Groth, R. Harris, K. Butler, J. Shields, S. Kharchenko and J. Douglas, Polymer, 2004, 45, 4227–4239. 17. B. Schartel, P. Poetschke, U. Knoll and M. Abdel-Goad, Eur. Polym. J., 2005, 41, 1061. 18. F. Gao, G. Beyer and Q. Yuan, Polym. Degrad. Stab., 2005, 89, 559. 19. S. Lohnmayer, Die speziellen Eigenschaften der Kunststoffe, Gafenau, Expert Verlag, 1984. 20. E. Ferrage, F. Martin, A. Boudet, S. Petit, G. Fourty, F. Jouffret, P. Micoud, P. De Parseval, S. Salvi, C. Bourgerette, J. Ferret, Y. SaintGerard, S. Buratto and J.P. Fortune, J Mater Sci, 2002, 37, 1561–1573. 21. S. Radhakrishnan, P. Sonawane and N. Pawaskar, J Appl Polym Sci, 2004, 93, 615–623. 22. L. Wang and J. Sheng, Polymer, 2005, 46, 6243–6249. 23. N. Grassie and G. Scott, Polymer degradation and stabilisation, Cambridge University Press, Cambridge, 1985. 24. A. Guyot, Polym. Degrad. Stab., 1986, 15, 219–235. 25. S. Ceccia, D. Ferri, D. Tabuani, P.L. Maffettone, Rheologica Acta, 2008, 47, 425–433.

CHAPTER 9

Combination of Carbon Nanotubes with Fire Retardants: Thermal and Fire Properties of Polystyrene Nanocomposites FLORENTINA TUTUNEA AND CHARLES A. WILKIE Department of Chemistry, Marquette University, PO Box 1881, Milwaukee, WI 53201, USA

9.1 Introduction Together with numerous advantages that synthetic polymeric materials provide to society in everyday life, there is one obvious disadvantage – the flammability of many synthetic polymers.1 Polymers are used in the manufacture of not only bulk parts, but also of films, fibres, coatings and foams. Fire hazard is a combination of factors, including ignitability, ease of extinction, flammability of the volatile products generated, amount of heat released on burning, rate of heat release, flame spread, smoke obscuration and smoke toxicity, as well as the fire scenario.2–5 Polymers are used in more and more applications and specific mechanical, thermal and electrical properties are required. Flame retardants may function in the gas phase, by removing the reactive hydrogen and hydroxyl radicals, or in the condensed phase, influencing the course of pyrolysis by forming char. The char insulates the polymer from the flame front and results in fewer

Fire Retardancy of Polymers: New Strategies and Mechanisms Edited by T Richard Hull and Baljinder K Kandola r Royal Society of Chemistry 2009 Published by the Royal Society of Chemistry, www.rsc.org

125

126

Chapter 9

combustible volatiles. The typical vapour-phase fire retardant is a halogen compound, most commonly an aromatic bromine compound. These function through the formation of HBr, which can react with and remove the hydrogen atoms and hydroxyl radicals which make up the flame. Among the nonreactive additives used as flame retardants is alumina trihydrate (ATH). Unlike the reactive halogen flame retardants, ATH reduces flammability by endothermically releasing water as it decomposes. Since this reaction is endothermic, it absorbs the heat of combustion, and thereby lowers the surface temperature near the resin. In addition, the evolved water dilutes the flammable gaseous reactants in the flame. The use of ATH also results in low smoke and toxic gas evolution.5 Phosphorus-containing flame retardants may function in either the vapour or condensed phase. In the vapour phase they act as radical traps similar to the process for a halogen. In the condensed phase, a phosphorus compound may change the degradation pathway of the polymer to promote char formation or the phosphorus compound may form a glassy-like layer which can act as a thermal barrier. A new class of materials, called nanocomposites, avoids one of the disadvantages of the traditional flame retardant systems, since typically the mechanical properties are enhanced. Generally the term ‘‘nanocomposites’’ describes a two-phase material with a suitable nanofiller dispersed in a polymer matrix at the nanometer (10 9) scale.6 Nanocomposites also show increased thermal stability and an improvement in flame retardancy at very low filler levels. The formation of char is responsible for these improved properties.7–9 Nanocomposite formation reduces the heat release rate (HRR), but they typically ignite more easily and all of the polymer will burn. For this reason, it is widely believed that nanocomposite formations combined with other materials may be an effective fire retardant system.10 Several papers from Kashiwagi and coworkers11–13 describe the effect of the addition of carbon nanotubes (CNTs) to polymers; the presence of CNT enhances the thermal stability of the polymer. The CNTs were at least as effective a flame retardant as organoclays and can be used at lower levels, since the char that is formed is denser. This chapter reports on the combination of the CNTs with decabromodiphenyl oxide (DECA), synergized with antimony oxide (Sb2O3), resorcinol diphosphate (RDP) and ATH. The combination of an organically modified clay along with the CNT and the flame retardants has also been examined.

9.2 Experimental 9.2.1 Materials The polymer used in this study was polystyrene (PS) with average molecular weight B230 000, and number average molecular weight B140 000, softening point 107 1C (Vicat, ASTM D 1525) and melt index 7.5 g/10 min (ASTM D

Combination of Carbon Nanotubes with Fire Retardants

127

1238, 200 1C/5 kg). The clay used was: Cloisite 20A (organically modified clay in which the surfactant is dimethyl dihalogenated tallow) and was kindly provided by Southern Clay Products. Multiwall CNTs were kindly provided by Nanocyl S.A., Belgium; DECA and ATH, Martinal OL 104 LE, were provided by Albemarle, while RDP was provided by Supresta Inc. Sb2O3 was kindly provided by Laurel Industries.

9.2.2 Preparation of Composites Nanocomposites were prepared by melt blending, using a Brabender Plasticorder at 190 1C for 10 minutes at 60 revolutions per minute (rpm) and the PS and additives were placed into the mixer at the same time. After blending, the mixture was removed and allowed to cool.

9.2.3 Instrumentation Thermogravimetric analysis (TGA) was performed on a SDT 2960 Simultaneous DTA-TGA unit from TA Instruments, under a constant nitrogen flow of 70 ml min 1; samples were run in duplicate and average values are reported. The experiments were performed at a rate ramp of 20 1C min 1 from 100 to 600 1C. The samples are first equilibrated at 100 1C then heated to 600 1C. Temperatures are considered accurate to 2 1C, while the char remaining at 600 1C is considered to be accurate to 3%. Cone calorimeter analyses were performed on an Atlas Cone 2 instrument using a cone-shaped heater, according to ASTM E 1354 at an incident flux of 35 kW m 2; the spark was continuous for 10 seconds after the sample ignited. All samples were run in triplicate and the average value is reported. The specimens for cone calorimetry were prepared by compression moulding of the sample (about 30 g) into 3  100  100 mm3 square plaques. Typical results from the cone calorimeter are reproducible to 10%, based on many thousands of samples that have been run.

9.3 Results and Discussion In this study representatives of three classes of fire retardants were used: bromine, DECA, synergized with Sb2O3 at a DECA:antimony ratio of 4:1; phosphorus, RDP; and minerals, ATH. The combination of the polymer with the flame retardant additive is considered first of all, and then combinations of the flame retardants and nanomaterials are examined.

9.3.1 Thermogravimetric Analysis The important parameters from the TGA curves are the onset temperature of the degradation, which is taken as the temperature at which 10% of the sample

128

Chapter 9

mass is lost, T10; the mid-point temperature of the degradation, T50; and the non-volatile residue at 600 1C, or char. Various loadings of ATH were melt blended with PS and the results are summarized in Table 9.1 and Figure 9.1. The temperature at which 10% mass loss occurs decreases as the amount of ATH increases, which may result from ATH decomposition, but the T50 increases, indicating enhanced thermal stability. The amount of residue is what is expected based on the ATH loading.

Table 9.1

TGA summary results for PS with fire retardants.

PS PS–15%RDP PS–30%RDP PS–7%ATH PS–14%ATH PS–21%ATH PS– 28%ATH PS–5%DECA–1.25%Sb2O3 PS–10%DECA–2.5%Sb2O3 PS–15%DECA–3.75%Sb2O3 PS–20%DECA–5%Sb2O3

Figure 9.1

T10 (1C)

T50 (1C)

Char (%)

402 391 391 396 388 388 381 382 379 376 377

426 428 436 427 429 435 439 412 406 402 401

0 1 2 5 10 14 19 2 3 5 6

TGA results for PS–ATH.

129

Combination of Carbon Nanotubes with Fire Retardants 14

Previous work in this laboratory was carried out on PS with RDP. These results are also presented in Table 9.1; the combination of RDP and PS leads to a decrease in the onset temperature and an increase in the mid-point temperatures, which may be due to the greater thermal stability of RDP. The combinations of DECA and antimony trioxide with PS are also reported in Table 9.1 and the TGA curves are given in the Figure 9.2; the onset temperature and T50 are decreased in the presence of the additives. To investigate possible synergy between nanodimensional materials and flame retardants with PS, the same loadings of fire retardants were used and an initial survey was performed at 1.5% CNT loading; the data are collected in Table 9.2 and the TGA curves are given in Figure 9.3. Except for the case of DECA, both T10 and T50 increase when both additives are present and the amount of char is greater than expected. Due to the rather considerable price of CNT, it was decided to reduce the CNT loading to 0.2 or

Figure 9.2

Table 9.2

TGA results for PS–DECA.

TGA summary results for PS–1.5%CNT–fire retardants.

PS PS–1.5%CNT PS–1.5%CNT–7%ATH PS–1.5%CNT–1.5%RDP PS–1.5%CNT–30%RDP PS–1.5%CNT–7%DECA

T10 (1C)

T50 (1C)

Char (%)

402 409 408 422 423 393

426 438 452 456 460 422

0 0 14 5 4 2

130

Chapter 9

Figure 9.3

TGA results for PS–CNT–fire retardants.

Figure 9.4

TGA results for PS–0.2%CNT and PS0.5%CNT.

0.5%. As can be seen in Figure 9.4, the addition of this low amount of CNT has essentially no effect on the TGA parameters for PS. The combination of CNT with DECA and Sb2O3, reported in Table 9.3 and Figure 9.5, shows no increase in thermal stability by TGA.

131

Combination of Carbon Nanotubes with Fire Retardants

Table 9.3

TGA results for PS–CNT–DECA.

PS PS–0.2%CNT PS+0.2%CNT–5%DECA–1.25%Sb2O3 PS–0.2%CNT–10%DECA–2.5 %Sb2O3 PS–0.2%CNT–15%DECA–3.75%Sb2O3 PS–0.5%CNT PS–0.5%CNT–5%DECA–1.25%Sb2O3 PS–0.5%CNT–10%DECA–2.5%Sb2O3 PS–0.5%CNT–20%DECA–5%Sb2O3

Figure 9.5

T10 (1C)

T50 (1C)

Char (%)

402 400 377 363 372 402 375 376 377

426 424 411 400 400 425 411 405 397

0 1 2 2 6 1 2 3 5

TGA results PS–CNT–DECA.

Figure 9.6 gives the TGA curves of PS and 0.2% CNT with various loadings of ATH; the TGA results are summarized in Table 9.4. Note that the temperature for 10% degradation does not depend on ATH concentration until the ATH reaches 35%, when it drops dramatically. This is likely due to the degradation of ATH. As was previously seen with flame retardants only, the 50% mass-loss temperature exhibits a considerable increase with an increase of the ATH loading. When 1.5% CNT was used, the amount of char was much greater than expected, but at the lower amounts of CNT, the char yields simply reflect the ATH loading. This suggests that there may be some interaction between CNT and ATH at high CNT loading, but not at low loading. The type of interaction has not yet been identified and this is being actively pursued.

132

Figure 9.6

Table 9.4

Chapter 9

TGA results for PS–CNT–ATH.

TGA results for PS–CNT–ATH. T10 (1C)

T50 (1C)

PS PS–0.2%CNT PS–0.5%CNT

402 400 402

426 423 425

0 0 0

PS–0.2%CNT–7%ATH PS–0.2%CNT–14%ATH PS–0.2%CNT–21%ATH PS–0.2%CNT–28%ATH PS–0.2%CNT–35%ATH

392 400 402 395 364

426 432 436 435 442

4 10 13 15 23

PS–0.5%CNT–7%ATH PS–0.%CNT–14%ATH PS–0.5%CNT–21%ATH PS–0.5%CNT–28%ATH

399 398 394 380

430 432 437 440

7 9 15 20

9.3.1.1

Char (%)

Addition of Clay to CNT-containing Systems

Beyer has claimed some synergy when CNT and clay are combined.15 Accordingly, an organically modified montmorillonite (Cloisite 20A) and CNT were combined with ATH and the results are given in Table 9.5 and Figures 9.8 to 9.12. The addition of 2% clay to PS gives no change in T10 but a significant increase in T50, shown in Table 9.5 and Figure 9.7. The addition of clay to the ATH-CNT system shows a significant increase in both T10 and T50, shown in Figures 9.8 to 9.11; Figure 9.12 provides a summary for all of amounts of ATH.

133

Combination of Carbon Nanotubes with Fire Retardants

Table 9.5 PS PS–2%clay

TGA results for PS–clay. T10 (1C)

T50 (1C)

Char (%)

402 404

426 444

0 0

Figure 9.7

TGA results for PS–clay.

Figure 9.8

TGA results for PS nanocomposites with clay and 7% ATH.

134

Chapter 9

Figure 9.9

TGA results for PS nanocomposites with clay and 14% ATH.

Figure 9.10

9.3.1.2

TGA results for PS nanocomposites with clay and 21% ATH.

Summary of the TGA Data

The addition of either 0.2 or 0.5% CNT to a PS–DECA or PS–ATH system has little effect on the TGA parameters. When 1.5% CNT is used, both T10 and T50 are increased relative to the use of the flame retardant alone. The combination of CNT and an organically modified clay with ATH leads to a significant enhancement in T50.

Combination of Carbon Nanotubes with Fire Retardants

Figure 9.11

TGA results for PS nanocomposites with clay and 28% ATH.

Figure 9.12

TGA results for PS with CNT, clay and various amounts of ATH.

135

9.3.2 Cone Calorimeter Evaluation The evaluation of the fire properties of the polymer systems uses the cone calorimeter. The parameters that are investigated with cone calorimetry are the:  time to ignition (TTI);  HRR, and especially its peak value (PHRR);

136

Chapter 9

 total heat release (THR);  specific extinction area (SEA), a measure of the amount of smoke produced during the combustion; and  mass-loss rate (MLR). Cone calorimetry provides important information on the fire properties of materials and it also provides useful information on nanocomposite formation. In work carried out in these laboratories, it has been observed that polymer– clay microcomposites do not show significant reduction in PHRR, but nanocomposites do show a significant reduction; this means that one can evaluate nanocomposite formation using the cone calorimeter.16–18 The amount of the reduction in the case of PS is up to 60%.19,20 As with the TGA work, we begin by showing the cone results for the combination of PS with only the fire retardants (Table 9.6). With the PS–RDP system, there is a 64% reduction in PHRR at 30% loading with a 55% reduction in the THR; the TTI also shows a small increase. In the ATH system, shown in Figure 9.13, the reduction in PHRR varies from 14 at 7% ATH to 53% at 28% ATH. There is not much change in the THR and the TTI was constant, except at 28% ATH where a large increase is seen. When ATH is present, these curves are similar to those reported by Schartel and Hull for thick samples, without charring at low amounts of ATH and with charring at high amounts.21 For DECA–Sb2O3, a reduction of 64% was obtained at 25% additive loading (20%DECA–5%Sb2O3), the PHRR is about the same at 15% DECA and it is lower at both 5% (26%) and 10% (43%) DECA.

Table 9.6

Cone calorimetry summary of PS–fire retardant systems.

Formulation

PHRRa (kW m 2) (% reduction)

THRb (MJ m 2)

ASEAc (m2 kg 1)

AMLRd (g s 1 m 2)

TTIe(s)

PS PS–15%RDP PS–30%RDP

1160  43 592  30 (49) 420  30 (64)

91  7 57  5 41  3

1330 1551 1852

27.4  1.1 15  1 14  1

60  0 63  1 77  6

PS–ATH7% PS–ATH14% PS–ATH21% PS–ATH28%

1002  102 (14) 787  18 (32) 608  21 (48) 550  24 (53)

90  4 93  1 89  3 84  2

1240  25 1214  12 1243  79 1309  98

23.5  0.6 20.1  0.7 13.2  0.7 9.5  0.1

37  1 37  4 37  2 51  2

821  25 (26) 662  229 (43) 487  63 (61)

52  7 34  3 27  2

1705  74 2333  74 2178  23

25.7  6.4 29.3  2.6 32.9  2.8

45  6 57  2 55  2

422  20 (64)

30  0

2127  36

33.7  1.7

55  6

PS–5%DECA–1.25%Sb2O3 PS–10% DECA–2.5%Sb2O3 PS–15%DECA– 3.75%Sb2O3 PS–20%DECA–5%Sb2O3 a

PHRR, peak heat release rate. THR, total heat released. ASEA, average specific extension area. d AMLR, average mass-loss rate. e TTI, time to ignition. b c

Combination of Carbon Nanotubes with Fire Retardants

Figure 9.13

HRR curves for PS–ATH.

Figure 9.14

HRR curves for PS–DECA.

137

The reduction in THR is 43% at the lowest DECA loading (5%) and 67% at the highest loading (20%), while the TTI shows no significant increase. The HRR curves are shown in Figure 9.14, and are similar to those of thick, noncharring samples when both DECA and Sb2O3 are present.

138

Figure 9.15

Table 9.7

Chapter 9

HRR results for PS–nanocomposites.

Cone results for PS–CNT–fire retardants.

Formulation

PHRRa (kW m 2) (% reduction)

THRb (MJ m 2)

ASEAc (m2 kg 1)

AMLRd (g/s 1 m 2) TTIe(s)

PS PS–1.5%CNT–15%RDP PS–1.5%CNT–30%RDP PS–1.5%CNT–7%ATH PS–1.5%CNT–7%DECA PS–1.5%CNT

1362 504  16 (63) 419  14 (69) 422  20 (69) 732  64 (46) 728  222 (47)

109 67  3 53  1 77  9 70  1 74  60

1507 2311  44 2483  161 1523  76 1904  105 1411  250

29.1 20.7  0.5 21.2  1.7 10.4  0.5 25.7  0.5 18.5  4.8

53 51  3 52  1 31  3 41  4 37  2

a

PHRR, peak heat release rate. THR, total heat released. c ASEA, average specific extension area. d AMLR, average mass-loss rate. e TTI, time to ignition. b

The cone results for the fire retardant systems that contain 1.5% CNT are shown in Figure 9.15 and the data are summarized in Table 9.7. The addition of 1.5% CNT to PS brings about a large reduction in PHRR, and the curves become much flatter, as well as a large reduction in THR. The reduction in THR is surprising since there is no change upon the addition of clay. The reduction in PHRR is much larger in the presence of 1.5% CNT than it is for the additive alone, except for the case of 30% RDP. When the amount of CNT used was 0.2 or 0.5%, there was no effect on the PHRR of PS, unlike the situation for 1.5% CNT; this is shown in Figure 9.16.

Combination of Carbon Nanotubes with Fire Retardants

Figure 9.16

139

HRR results for PS with 0.2 and 0.5%CNT.

This plot also shows that the dramatic decrease in the TTI at 0.2% loading (which perhaps results from an increase in melt viscosity) is independent of the reduction in PHRR, which is often ascribed to the formation of a barrier. This is a suitable topic for further investigation because it implies that there must be some interaction between the larger amount of CNT and the flame retardants, as noted earlier from TGA data. The HRR plots for the additives plus clay are given in Figure 9.17. The shape of the curves shows a drastic change compared to PS or PS–CNT. The presence of the additives makes the sample thermally thicker. The cone data of the PS–CNT–DECA system are gathered and reported in Table 9.8. For DECA alone at 10%, there is a 43% reduction, while when either 0.2 or 0.5% CNT is added, the reduction in PHRR is better than that seen with only 20% DECA – 69%. This must be contrasted to the case where 1.5% CNT was added and there was no change in the PHRR. The results when various amounts of ATH and 0.2 and 0.5% CNT are combined are shown in Table 9.9 and Figure 9.18. The reduction in PHRR is about the same for each ATH loading at both 0.2 and 0.5% CNT. The addition of the CNT causes a decrease in the PHRR, which varies between 5 and 8%. This is within the typically quoted error bars on cone measurements, so there may be no change; most likely there is no additional effect on the combination of ATH and CNT with PS. TTIs for both 0.2 and 0.5% CNT loadings decrease compared with the pure polymer, but this decrease is not a function of loading. THR is unchanged, except for those systems which contain CNT and the highest loading of ATH. Since the THR is unaffected by the addition of either

140

Chapter 9

Figure 9.17

Table 9.8

HRR results for PS nanocomposites containing DECA.

Cone results for PS–CNT–DECA.

Formulation PS PS–5%DECA– 1.25%Sb2O3 PS–5%DECA– 1.25%Sb2O3– 0.2%CNT PS–5%DECA– 1.25%Sb2O3– 0.5%CNT PS–10%DECA– 2.5%Sb2O3 PS–10%DECA– 2.5%Sb2O3– 0.5%CNT PS–20%DECA– 5%Sb2O3 PS–20%DECA– 5%Sb2O3– 0.5%CNT a

PHRRa (kW m 2) (% reduction)

THRb (MJ m 2)

ASEAc (m2 kg 1)

AMLRd (g s 1 m 2)

TTIe (s)

1208  127 821  25 (26)

105  4 52  7

1362  33 1705  74

27.7  2.1 25.7  6.4

60  1 45  6

646  8 (42)

46  3

2063  40

30.3  1.7

42  3

596  95 (47)

49  4

1959  102

29.3  1.4

37  1

662  229 (43)

34  3

2333  74

29.3  2.6

57  2

380  22 (69)

31  0

2608  45

28.7  0.6

45  4

422  20 (64)

30  0

2127  36

33.7  1.7

55  6

394  23 (67)

24  13

1961  126

28.3  2.9

41  2

PHRR, peak heat release rate. THR, total heat released. c ASEA, average specific extension area. d AMLR, average mass-loss rate. e TTI, time to ignition. b

141

Combination of Carbon Nanotubes with Fire Retardants

Table 9.9

Cone results for PS–CNT–ATH.

Formulation

PHRRa (kW m 2) (% reduction)

PS PS–0.2%CNT PS–0.5%CNT PS–2%clay

1076  41 1110  65 (0) 1013  45 (9) 971  35 (10)

PS PS–0.2%CNT–7%ATH PS–0.2%CNT–14%ATH PS–0.2%CNT–21%ATH PS–0.2%CNT–28%ATH PS–0.2%CNT–35%ATH

1233  93 926  44 761  34 579  33 486  37 397  22

PS PS–CNT 0.5%–ATH7% PS–CNT0.5%–ATH14% PS–CNT. 5%–ATH21% PS–CNT0.5%–ATH28%

THRb (MJ m 2)

ASEAe (m2 kg 1)

AMLRd (g s 1 m 2)

TTIe (s)

88  3 96  2 96  1 94  2

1113  36 1197  19 1239  11 1232  17

29.4  1.2 27.3  1.7 25.6  0.6 25.6  0.6

56  8 41  1 39  2 59  3

(25) (38) (53) (61) (68)

91  2 92  3 90  1 83  4 82  4 73  0

1122  60 1227  21 1259  25 1233  24 1205  46 1106  37

31.5  2.7 25.0  0.8 18.8  0.9 13.3  0.2 11.3  0.8 6.6  0.3

53  2 28  3 32  2 31  3 31  0 38  2

1208  127 798  20 (34) 600  16 (50) 516  22 (57) 459  20 (62)

105  4 105  3 100  0 94  5 84  1

1362  33 1422  38 1363  10 1342  29 1338  40

27.7  2.1 19.6  0.2 13.2  0.3 10.4  0.1 8.4  0.2

60  1 29  3 31  2 36  1 38  5

a

PHRR, peak heat release rate. THR, total heat released. ASEA, average specific extension area. d AMLR, average mass-loss rate. e TTI, time to ignition. b c

Figure 9.18

HRR results for PS–ATH–0.2%CNT.

142

Chapter 9

of these separately, this suggests that some synergistic or additive effect may be occurring; this requires further investigation.

9.3.2.1

Combination of Clay and CNT

The cone calorimeter results for the combinations of clay with CNT and ATH are shown in Table 9.10 and the heat-release plots are provided in Figures 9.19 to 9.23. For comparison, the PHRR decreases by about 50% upon the addition of 3% of this clay.22 At low ATH loadings, especially 7% ATH, there is a very large decrease in the PHRR upon the addition of the clay. As the ATH loading increases, the values of the PHRR come closer together so that at 21 or 28% ATH loading, there is essentially no difference in the PHRR in the presence or absence of the clay.

9.3.2.2

Summary of Cone Calorimetric Data

The combination of 10% DECA and 0.5% CNT gives the same value for the PHRR as is obtained when only 20% DECA is used so clearly this is an Table 9.10

Cone results for PS–fire retardants–nanocomposites. PHRRa (kW m 2) (% reduction)

THRb (MJ m 2)

ASEAc (m2 kg 1)

AMLRd (g s 1 m 2)

TTIe (s)

PS PS–0.2%CNT–7%ATH– 2%clay PS–.2%CNT0–14%ATH– 2%clay PS–0.2%CNT–21%ATH– 2%clay PS–0.2%CNT–28%ATH– 2%clay PS–0.2%CNT–35%ATH– 2%clay

1076  41 579  41 (46)

88  3 89  2

1113  36 1387  13

29.4  1.2 16.2  0.7

56  8 29  4

506  15 (53)

88  1

1353  10

11.3  0.2

31  1

397  12 (63)

81  1

1317  66

7.8  0.2

31  1

328  8 (69)

75  1

1229  48

6.1  0.2

36  2

309  12 (71)

65  0

1159  13

6.0  0.1

37  2

PS–0.5%CNT–7%ATH– 2%clay PS–0.5%CNT–14%ATH– 2%clay PS–0.5%CNT–21%ATH– 2%clay PS–0.5%CNT–28%ATH– 2%clay PS–0.5%CNT–35%ATH– 2%clay

591  14 (45)

91  1

1391  15

15.9  0

30  2

494  9 (54)

86  0

1413  11

10.9  0.2

30  1

382  9 (65)

79  2

1362  61

5.3  4.6

31  1

352  17 (67)

73  0

1242  35

7.1  0.3

37  1

317  21 (71)

68  1

1071  58

5.9  0.3

42  1

Formulation

a

PHRR, peak heat release rate. THR, total heat released. ASEA, average specific extension area. d AMLR, average mass loss rate. e TTI, time to ignition. b c

Combination of Carbon Nanotubes with Fire Retardants

Figure 9.19

HRR results for PS–fire retardants–nanocomposites.

Figure 9.20

HRR results for PS–fire retardants–nanocomposites.

143

advantageous system since it permits one to significantly decrease the halogen loading. It is surprising that at 1.5% CNT there is no change in the PHRR because of the presence of the clay. This suggests that the amount of CNT is very important. In contrast, with ATH at low loading, 0.2 or 0.5%, there does

144

Chapter 9

Figure 9.21

HRR results for PS–fire retardants–nanocomposites.

Figure 9.22

HRR results for PS–fire retardants–nanocomposites.

Combination of Carbon Nanotubes with Fire Retardants

Figure 9.23

145

HRR results for PS–fire retardants–nanocomposites.

not seem to be an effect, but at 1.5% there is a definite interaction between ATH and CNT. For this system, there does not appear to be any advantage in combining an organically modified clay with CNT when evaluated by cone calorimetry.

9.4 Conclusions The combination of CNTs at 0.2 or 0.5% with either ATH or RDP does not appear to give any significant enhancement in thermal stability as measured by TGA or in fire properties as measured by cone calorimetry compared to the use of the flame retardant alone. When 1.5% CNTs are used, there is apparently an interaction between ATH and the CNT which is manifested in a large amount of char and a greatly decreased value for the PHRR. Clearly some synergistic interaction occurs between CNT and DECA at the low, 0.2 or 0.5% CNT, loading and the reduction in the PHRR at 10% DECA–CNT is the same as that observed for 20% DECA alone. It is very clear that there is some potential advantage in combining CNT with fire retardants in some polymers. The challenge that faces us is to identify the materials which should be used, meaning the fire retardant and the polymer. In addition, the amount of material that should be added must also be ascertained through a trial-and-error process. In this work, it has been shown that there is some advantage in combining 1.5% CNT with ATH in PS. To be practical, the amount of CNT must be reduced and it has been shown that reducing it to 0.5% is too little. It may be that there are other polymers, for instance a

146

Chapter 9

polyolefin like polyethylene or polypropylene, that can be adequately protected by the addition of CNT with a fire retardant. Future work is required to determine this.

References 1. S.V. Levchik, in Flame retardant polymer nanocomposites, A.B. Morgan and C.A. Wilkie eds., Wiley Interscience, 2007, pp. 1–29. 2. M. Hirschler, in Fire and Polymers: Materials and Solutions for Hazard Prevention. Eds. G.L. Nelson and C.A. Wilkie, ACS Symposium Series 797, American Chemical Society, Washington, DC, 2001, pp. 293–306. 3. D.A. Purser, Polym. Int., 2000, 49, 1232. 4. D.J. Irvine, J.A. McCluskey and I.M. Robinson, Polym. Degrad. Stab., 2000, 67, 383. 5. S. Salman and D. Klempner, Plast. Eng., 1979, 35, 39. 6. G. Beyer, Polym. Adv. Tech., 2006, 17, 218. 7. G. Beyer, Polym. News, 2001, 26, 370. 8. M. Le Bras, G. Camino, S. Bourbigot and R. Delobel eds., Fire Retardancy of Polymers: The use of intumescence., Royal society of Chemistry, Cambridge, 1998. 9. J.W. Gilman, T. Kashiwagi and J. Lichtenhan, SAMPE J, 1997, 4, 40. 10. Y. Hu and L. Song, in Flame retardant polymer nanocomposites, A.B. Morgan and C.A. Wilkie eds., Wiley Interscience, 2007, pp. 191–233. 11. T. Kashiwagi, E. Grulke, J. Hilding, R. Harris, W. Awad and J. Douglas, Macromol. Rapid Comm., 2002, 23, 761. 12. T. Kashiwagi, E. Grulke, J. Hilding, K. Grith, R. Harris, K. Butler, J. Shields, S. Kharchenko and J. Douglas, Polymer, 2004, 45, 4227. 13. T. Kashiwagi, F. Du, K. Winey, K. Groth, J. Shields, S. Bellayer, H. Kim and J. Douglas, Polymer, 2005, 46, 471. 14. G. Chigwada and C.A. Wilkie, Polym. Degrad. Stab., 2003, 80, 551. 15. G. Beyer, Fire Mater., 2005, 29, 61. 16. S. Su, D.D. Jiang and C.A. Wilkie, Polym. Degrad. Stab., 2004, 83, 301. 17. S. Su, D.D. Jiang and C.A. Wilkie, Polym. Degrad. Stab., 2004, 83, 321. 18. S. Su, D.D. Jiang and C.A. Wilkie, Polym. Adv. Tech., 2004, 15, 225. 19. J. Zhu, P. Start, K.A. Mauritz and C.A. Wilkie, Polym. Deg. Stab., 2002, 77, 253. 20. M. Zanetti, G. Camino, D. Canavese, A.B. Morgan, F.J. Lamelas and C.A. Wilkie, Chem. Mater., 2002, 14, 189. 21. B. Schartel and T.R. Hull, Fire Mater., 2007, 31, 327. 22. J. Zhu and C.A. Wilkie, Polym. Int., 2000, 49, 1158.

CHAPTER 10

Significant Assessment of Nanocomposites’ Combustion Behaviour by the Appropriate Use of the Cone Calorimeter A. FINA, F. CANTA, A. CASTROVINCI1 AND G. CAMINO Centro di Cultura per l’Ingegneria delle Materie Plastiche – Politecnico di Torino, V.le T. Michel, 5-15100 Alessandria, Italy; 1 Current address: University of Applied Sciences of Southern Switzerland (SUPSI), Galleria 2, CH-6928 Manno, Switzerland

10.1 Introduction The cone calorimeter has become one of the most important tools for research and development of fire-retardant polymeric materials. Several parameters are obtained from this test,1 but the most often used is the peak of heat release rate (PHRR). However, a major problem arises from the simple comparison of PHRR values, which may be misleading when polymer materials to be compared are characterized by very different combustion behaviours. In particular, charring strongly affects the shape of the HRR curve. Char-forming materials show the PHHR at an early stage of the combustion because of the protective action of the char building up on the surface of the burning sample, whereas non-charring materials (including most of the thermoplastic polymers) show the PHHR at the end of the burning process.2–4

Fire Retardancy of Polymers: New Strategies and Mechanisms Edited by T Richard Hull and Baljinder K Kandola r Royal Society of Chemistry 2009 Published by the Royal Society of Chemistry, www.rsc.org

147

148

Chapter 10

This is often the case when comparing pure bulk polymers, which usually show limited or no charring, and their correspondent nanocomposites, which often show the char-forming behaviour. In the past, the observed reductions in PHHR (up to 70%) for nanocomposites were directly related to improvements in the fire resistance, which now appear to be largely overestimated. This is also because combustion cone-calorimeter tests represent only a single and peculiar fire scenario, namely a well-ventilated forced combustion condition.4 The PHHR at the end of the burning phase for non-charring polymers is caused by the thermal feedback that occurs when the pyrolysis zone reaches the rear surface of the sample. During combustion, part of the heat is absorbed by the underlying material, but when the pyrolysis front approaches the base of the specimen bottom, heat is no longer removed by any remaining underlying material, which results in an increased material-heating rate. This results in a faster pyrolysis and subsequent increase of the measured HRR. This particular behaviour is determined by the standard insulated sample holder set-up in the cone calorimeter. Indeed, the large PHHR at the end of burning of non-charring materials has been demonstrated to disappear using a modified sample holder that conducts the heat away from the back of the sample by a ‘‘heat sink’’ effect.4,5 However, PHRR values for charring materials are observed just after ignition in the earlier stages of the protective char-layer build-up and usually the shape of the HRR curve is not significantly affected by thermal feedback, so that experimental set-up plays a minor role on the PHRR for these materials. Based on these observations, to evaluate fire performance in terms of PHRR for charring materials in comparison with non-charring polymer is certainly problematic. In this chapter, a method to assess the actual fire performance of nanocomposites is proposed in terms of a ‘‘steady value of HHR’’, defined as the constant rate of combustion observed in the time range between ignition and the occurrence of a thermal feedback contribution.

10.2 Experimental 10.2.1

Materials

A commercial grade of polyamide 6 from Rhodia, PA6 Technyl C206, was used (density 1.14 g cm
3, melting point 222 1C). The nanofiller used is Cloisite 30B from Southern Clay Products Inc, a natural montmorillonite modified by ionexchange with a methyl, tallow, bis(2-hydroxyethyl), quaternary ammonium chloride. The choice of components used was based on results published in the literature which show optimal nanoclay dispersions into a PA6 matrix and the formation of an exfoliated nanocomposite structure.6–9

10.2.2

Preparation and Characterization

The nanocomposite was prepared by melt processing using a Leistritz 27 co-rotating twin-screw extruder (d ¼ 27 mm, l/d ¼ 40) equipped with a low

149

Significant Assessment of Nanocomposites’ Combustion Behaviour

Temperature profile

Figure 10.1

Screw and temperature profiles used in compounding.

shear-stress screw profile (Figure 10.1), with 6 kg h
1 throughput and 100 revolutions per minute (rpm) screw speed. Polymer was dried for 4 hours at 80 1C and nanoclay for 4 hours at 100 1C to eliminate moisture before processing. Nanoclay was added at 5% weight percent (wt%), by a dedicated sidefeeder at 14 diameters length. Unfilled polymer was processed under the same conditions as the composites and a reference material. Plate specimens (100  100  8 mm) for the cone calorimeter tests were prepared by injection moulding using a Demag INT Elect 100/400 hydraulic press with the two heated plates at 60 1C using a maintenance pressure of 25 kPa (250 bar) for a cycle time of 61 s. Compounds were dried for 4 hours at 100 1C before processing to eliminate moisture. As the available injection moulding apparatus does not allow the production of specimens thicker than 8 mm, pairs of 8 mm specimens were laid one upon the other to obtain 16 mm thick samples. Improved thermal contact was achieved by gluing the two specimens together with a very small amount of cyanoacrylate glue. Composition and abbreviations for samples are given in Table 10.1 in which P stands for unfilled polymer, C for nanocomposite and the two associated figures stand for sample thickness and heat flux, respectively.

Table 10.1

Samples and imposed external heat fluxes.

Material PA6

Thickness (mm) 8 16

PA6+5 wt% Cloisite 30B

8 16

Heat flux (kW m
2)

Abbreviation

50 75 50 75 50 75 50 75

P8
50 P8
75 P16
50 P16
75 C8
50 C8
75 C16
50 C16
75

150

Chapter 10

X-Ray diffraction (XRD) patterns were obtained on a Thermo ARL XTRA48 diffractometer using Cu Ka radiation (l ¼ 1.54062 A˚) and used to evaluate the nanoclay dispersion.

10.2.3

Combustion Tests

Combustion tests were carried out on a Fire Testing Technology (FTT) cone calorimeter apparatus. The tests were performed in the horizontal position without using the retainer frame and therefore surface-dependent data were calculated using the entire specimen area (100 cm2). Specimens were wrapped in aluminium foil leaving an upper edge of 3 mm and placed on ceramic backing boards at a distance of 25 mm from the cone base (Figure 10.2). This non-standard experimental set-up was adopted to prevent dripping of molten material, which is particularly significant for the pure polymer. External heat flux was set at either 50 or 75 kW m
2. These values are comparable to conditions of medium- and well-developed fire scenarios, respectively,4,10 and lead to reasonable burning times for the largest specimen thicknesses used; lower heat fluxes would result in very time-consuming experiments. All the tests were performed at least in duplicate and the average values are reported herein, along with their deviations with respect to the average. Irradiances were measured by placing the Schmidt-Boelter heat-flux meter (Medtherm corp, mod. GWT-10-32-485A) of the cone calorimeter equipment at various distances from the cone base, in the centre of specimen position.

Figure 10.2

Specimen mounting set-up.

151

Significant Assessment of Nanocomposites’ Combustion Behaviour 6000

5000

PA6 + 5wt.% Cloisite 30B Cloisite 30B

Intensity [cps]

4000

3000

2000

1000

0 2

Figure 10.3

4

6

8 2θ angle [°]

10

12

14

XRD patterns for Cloisite 30B and PA6+Cloisite 30B nanocomposite.

10.3 Results and Discussion PA6 nanocomposite morphology was evaluated by XRD, which showed no diffraction peaks related to the nanoclay interlayer distance (Figure 10.3), suggesting exfoliation of nanoclay platelets in the PA6 matrix. This is in agreement with results from multitechnique characterization [XRD, scanning electron microscopy (SEM), transmission electron microscopy (TEM), rheology, nuclear magnetic resonance (NMR)] reported elsewhere9 on an equivalent PA6 nanocomposite prepared in the same experimental conditions. Heat-release parameters obtained from cone calorimeter analyses on 8 and 16 mm specimens are discussed below, including PHRR, average heat release rate (AHRR, the average value between time to ignition and end of test, obtained with ISO5660 criterion), and total heat evolved (THE), as well as time-to-ignition (TTI) and mass loss. The discussion of other parameters obtained from the same analyses, such as smoke, CO and CO2 production, is outside the objective of this work.

10.3.1

8 mm Specimens Combustion Behaviour

Results obtained from P-8 mm and C-8 mm samples are summarized in Table 10.2. A reduction of TTI is generally observed with increasing heat flux, because of the higher heating rate on the specimen surface, e.g. for neat PA6, TTI is 84 s at 50 kW m
2 and 38 s at 75 kW m
2.

152

Table 10.2

P8
50 P8
75 C8
50 C8
75

Chapter 10

Cone calorimeter results for 8 mm thick sample. TTI (s)

Mass loss (%)

PHRR (kW m
2)

AHRR (kW m
2)

THE (MJ m
2)

84  6 38  1 54  4 30  3

100  1 100  1 97  1 97  1

1123  56 1466  124 480  14 663  24

604  18 745  19 364  20 504  17

263  1 263  2 245  1 244  4

At both heat fluxes, a significant reduction of TTI for the nanocomposite is observed as compared with that of pure PA6, with ignition at 54 s for 50 kW m
2 and at 30 s for 75 kW m
2. Both lower11,12 and higher13,14 TTI values have been previously reported in the literature for layered silicate PA6 nanocomposites compared with reference PA6 and generally show little agreement. From a general point of view, a variable trend for TTI is observed for nanocomposites compared to pure polymers, which may increase or decrease depending on the type of polymer and nanofiller, as well as on processing and testing conditions.15 The anticipated ignition observed in this work could be ascribed to the thermal instability of the quaternary alkyl ammonium nanofiller organomodifier and/or to the physical effects of nanoclay on the specimens’ surfaces. Indeed, nanoclay platelets may affect the flow of molten material prior to ignition by the well-known increase in the molten viscosity,9,16–18 thereby modifying thermal convection flow and resulting in a higher surface-heating rate. However, the mechanistic study of ignition in PA6 nanocomposites requires further studies beyond the objective of this work. Comparison between pure PA6 and nanocomposite in terms of THE shows a limited reduction for the nanocomposite at both heat fluxes, which is only slightly higher than that calculated on the basis of the nanoclay inorganic content. This means that nanoclay is not effective at preventing PA6 from combustion; nevertheless, strong effects on the HHR are induced by the nanoclay in PA6. HHR curves for 8 mm specimens are reported in Figure 10.4. Pure PA6 shows the typical plot for a non-charring polymer, at both 50 and 75 kW m
2, evidenced by the peak at the end of the burning process. Despite the curve being strongly asymmetric, no HRR plateau is observed, showing an intermediate behaviour between the typical thermally thin and thermally thick behaviours.2,4 In contrast, the nanocomposite shows a quasi steady-state combustion at both 50 and 75 kW m
2, with plateaus at ca. 440 kW m
2 and 630 kW m
2, respectively. It is furthermore clearly observable that specimens burned under the higher external heat flux (75 kW/m2) show higher HHRs over all the well-sustained combustion time range as compared with the corresponding material tested at 50 kW/m2. At the lower heat flux, the PHRR values are 1123 kW/m2 for PA6 and 480 kW/m2 for the nanocomposite; whereas at 75 kW/m2 values are increased to 1466 kW/m2 and 663 kW/m2, respectively. Based on the PHRR reductions (ca. 55% at both heat fluxes) between pure polymer and nanocomposites, one

153

Significant Assessment of Nanocomposites’ Combustion Behaviour 1600 P8-50 1400

HRR [kW/m2]

C8-50 1200

P8-75

1000

C8-75

800 630 kW/m2 600 440 kW/m2 400 200 0 0

100

200

300

400

500

600

700

800

Time [s]

Figure 10.4 Cone calorimeter HRR curves for 8 mm thick specimens at 50 and 75 kW m
2.

could predict a large improvement in fire performance for the nanocomposite compared with the pure PA6. However, the different contributions of thermal feedback in the two different materials mean that simple peak comparison is certainly misleading. The average HRR value could be considered a slightly more reliable parameter, reducing the effect of the error brought by this final peaking effect. Based on average HRR values reported in Table 10.2, a significant reduction in combustion rate is still evident when comparing pure polymer with its nanocomposite analogue, amounting to ca. 40% at 50 kW m
2 and ca. 32% at 75 kW m
2. The use of the HRR at the curve shoulder observed in intermediate-thickness specimens could, in principle, be used; however, the determination of an accurate value appears to be subjective in most cases, including the case of 8 mm thick PA6 plots here discussed. The proposal in this chapter is to evaluate the effectiveness of nanoclay in reducing HHR when transient contributions unrepresentative of the material burning are eliminated. To reach this condition, specimens with higher thickness were tested.

10.3.2

16 mm Specimens Combustion Behaviour

Results obtained from analyses of results from P-16 mm and C-16 mm samples at different heat fluxes are summarized in Table 10.3.

154

Chapter 10

Cone calorimeter results for 16 mm thickness sample.

Table 10.3

P16
50 P16
75 C16
50 C16
75

TTI (s)

Mass loss (%)

pkHRR (kW m
2)

avHRR (kW m
2)

Plateau HRR (kW m
2)

THE (MJ m
2)

71  3 28  1 60  2 21  1

100  1 100  1 97  1 96  1

870  9 1192  146 395  10 605  7

557  8 781  25 291  24 419  17

580  40 940  30



502  3 529  3 506  30 484  10

1200

P16-50

1000

940 kW/m2

C16-50 P16-75 C16-75

HRR [kW/m2]

800 605 kW/m2 580 kW/m2

600

400 395 kW/m2

200

0 0

Figure 10.5

500

1000 Time [s]

1500

2000

Cone calorimeter HRR curves for 16 mm thick samples at 50 and 75 kW m
2.

The effect of the nanofiller on TTI at both 50 and 75 kW m
2 follows the same trend observed for 8 mm specimens. At 50 kW/ m
2 the pure PA6 ignites at 71 s and the composite at 60 s; at 75 kW m
2 pure PA6 ignites at 28 s and the composite at 21 s. However, a general reduction of TTI in these thicker samples is observed, though no significant differences are expected when passing from intermediate thickness specimens to thermally thick ones in terms of ignition time.19 As expected, the total heat released for 16 mm specimens is double that of the correspondent 8 mm tests. The main features of the HRR plots obtained with 16 mm specimens (Figure 10.5) are the different shapes with respect of those obtained with 8 mm specimens, both for neat PA6 and the nanocomposite. P16
50 shows an HRR plateau at ca. 580 kW m
2, followed by the wellknown peak (870 kW m
2) at the end of combustion, whereas values for P16
75

155

Significant Assessment of Nanocomposites’ Combustion Behaviour
2


2

are ca. 940 kW m for the plateau and 1192 kW m for PHRR. Note that plateau values are not perfectly defined, but show a drop in the HRR curve at about half the time between ignition and flameout, probably because of the imperfect thermal contact between the two 8 mm specimens laid one upon the other. This observation shows that a single solid block is to be preferred, when such a high specimen thickness can be obtained. PHRR values observed with 16 mm PA6 specimens are significantly lower than those for the 8 mm specimens of the same material, in agreement with data previously reported for PMMA,4 possibly because of the lower thermal feedback irradiation at the end of 16 mm specimens burning due to the larger distance from the radiation source. However, measurements of irradiated heat flux from the cone heater coil show a limited decrease of incident irradiation with increasing distance from the cone base, as shown in Figure 10.6. Indeed, at (25+16) mm from the cone base, reductions are only of 3.6 kW m
2 at 50 kW m
2 on the specimen surface (25 mm from the cone base) and 4.2 kW m
2 at 75 kW m
2, this being in good agreement with results reported in literature for a lower imposed heat flux.1,4,5 The curves obtained for 16 mm nanocomposites also show other interesting features. Indeed, both at high and low heat flux, the HRR peak (395 kW m
2 for C16
50 and 605 kW m
2 for C16
75) is readily achieved after ignition, and approximately corresponds to the plateau previously observed with 8 mm specimens. However, after this first stage, a decrease in HRR is observed, because of the continuous accumulation of the nanoclay as a consequence of

75

Heat Flux [kW/m2]

70 65 60 55 50 45 40 24

Figure 10.6

26

28

34 38 36 30 32 Distance from Cone base [mm]

40

42

Decreases in imposed heat flux with distance from cone base.

156

Chapter 10

1 cm

Figure 10.7

Nanocomposite residue after combustion (top view).

polymer degradation, which produces a structured char able to slow down the feed of volatiles to the flame by the well-known barrier effect.20,21 As stated above, the PA6 nanocomposite used in this work does not lead to the formation of a continuous char layer on the specimen surface. Also, the accumulation of nanoclay in the earlier combustion stage is concentrated in spots distributed on the surface, which leads to the formation of column-like structures during combustion development (Figure 10.7). This is in good agreement with the formation of carbonaceous floccules reported by Kashiwagi et al. through the effect of convection flow and bubbling in the molten state during combustion.13 As a result of the accumulation of the inorganic nanofiller because of polymer ablation it is impossible to obtain a proper steady state of combustion for these polymer-layered nanocomposites. However, the value of the HRR peak observed just after ignition is representative of the material behaviour, and provides a meaningful index of char efficiency from the earliest stage of combustion. The results obtained for 16 mm specimens provide evidence of steady-state burning for the pure PA6. Therefore a proper comparison between PA6 and its corresponding nanocomposite in terms of HHR can be undertaken to evaluate material burning independently of the transient phenomena that occur just prior to flameout. An accurate and reliable value for the HRR reduction between pure PA6 and PA6 nanocomposite can therefore be calculated as: SteadyHRRPA6
PeakHRRPA6
nanocomposite  100 SteadyHRRPA6

ð10:1Þ

157

Significant Assessment of Nanocomposites’ Combustion Behaviour

Table 10.4

HRR reductions calculated on pkHRR reduction and with the steady state method (16 mm specimens).

Heat flux (kW m
2) PeakHRRPA6
PeakHRRPA6
nanocomposite PeakHRRPA6

 100

SteadyHRRPA6
PeakHRRPA6
nanocomposite SteadyHRRPA6

 100

50

75

55%

49%

32%

36%

Using this approach, the HRR reductions for the nanocomposite compared with pure PA6 amount to about 33% at both heat fluxes used in this work. For PHRR, reductions of about 50% would have been obtained, evidence of the overestimation of HHR reduction (Table 10.4).

10.4 Conclusions Based on the experimental results reported in this chapter, the ability to evaluate properly the combustion behaviour of a polymer nanocomposite during a forced combustion test has been shown. The correct comparison between the performances of the nanocomposite and of the reference polymer, independent of transient phenomena, may be obtained by evaluating the steadystate burning rate for the reference polymer, observable as a plateau on the HRR curve. We propose to increase the specimen’s thickness to achieve steady-state burning, a method that leads to a steady-state combustion range which is not affected by external effects, such as the ‘‘heat sink’’ effect induced by other nonstandard specimen mountings, for example the use of a conductive block underneath the specimen.4,5 With the approach described here, the effectiveness of 5% nanoclay in the reduction of the HHR of PA6 was assessed and quantified to be about 33% at both 50 kW m
2 and 75 kW m
2 incident flux levels, whereas about 50% reductions are observed for PHRR values. This shows the overestimation obtained when the transient phenomena that occur in non-charring materials during the final stage of combustion in standard cone calorimeter tests are not taken into account. The proposed method to calculate the HRR reduction is not intended to provide a single parameter with which to evaluate combustion behaviour. The comprehensive assessment of nanoclay effectiveness as a fire retardant must consider other parameters from the cone calorimeter test (such as TTI, THE, effective heat of combustion, smoke and CO productions) as well as performance in standard flammability tests [e.g. UL94 test, limiting oxygen index (LOI) etc.] and other tests specific for the final application of the nanostructured material. Note also that the effectiveness in HRR reduction calculated and discussed here may not directly provide a prediction of the actual fire performance in real material applications. Specimen configuration (e.g. vertical

158

Chapter 10

or horizontal orientation, back-insulation) and geometry (e.g. thickness, curvature) in real applications can be significantly different from those used in the cone calorimeter tests reported here. Nevertheless, we believe that the approach discussed in this chapter is useful for preliminary research on materials, particularly when details of end-user applications are not available.

Acknowledgements This study was carried out in the Frame STRP European research program ‘‘PredFire Nano’’, n1. 013998, in the sixth Framework Program.

References 1. V. Babrauskas, in Heat Releases in Fires, ed. V. Babrauskas and S. Grayson, Elsevier Applied Science, London, 1992, p. 61. 2. R. E. Lyon, M. L. Janssens, Polymer flammability: U.S. Department of Transportation Federal Aviation Administration technical report. 1982 Technical Report DOT/FAA/AR05/14, FAA, 2005. available online at http://www.fire.tc.faa.gov. 3. V. Babrauskas, Fire Mater., 1984, 8, 81. 4. B. Schartel and T.R. Hull, Fire Mater., 2007, 31, 327. 5. B. Schartel, M. Bartholmai and U. Knoll, Polym. Degr. Stab., 2005, 88, 540. 6. T.D. Fornes, P.J. Yoon, D.L. Hunter, H. Keskkula and D.R. Paul, Polymer, 2002, 43, 5915. 7. O. Monticelli, Z. Musina, A. Frache, F. Bellucci, G. Camino and S. Russo, Polym. Degr. Stab., 2007, 92, 370. 8. H.R. Dennis, D.L. Hunter, D. Chang, S. Kim, J.L. White, J.W. Cho and D.R. Paul, Polymer, 2001, 42, 9513. 9. F. Samyn, S. Bourbigot, C. Jama, S. Bellayer, S. Nazare, T.R. Hull, A. Castrovinci, A. Fina and G. Camino, Eur. Polym. J., 2008, 44, 1642. 10. V. Babrauskas, Fire Mater., 1995, 19, 243. 11. J.W. Gilman, Appl. Clay Sci., 1999, 15, 31. 12. S. Bourbigot, E. Devaux and X. Flambard, Polym. Degr. Stab., 2002, 75, 397. 13. T. Kashiwagi, R.H. Harris, X. Zhang, R.M. Briber, B.H. Cipriano, S.R. Raghavan, W.H. Awad and J.R. Shields, Polymer, 2004, 45, 881. 14. B.N. Jang and C.A. Wilkie, Polymer, 2005, 46, 3264. 15. A. Castrovinci and G. Camino, in Multifunctional Barriers for Flexible Structure: Textile, Paper and Leather, S. Duquesne, C. Magniez, G. Camino ed., Springer Verlag, Berlin, 2007, p. 87. 16. R. Krishnamoorti and K. Yurekli, Curr. Opin. Colloid Interface Sci., 2001, 6, 464. 17. L. Incarnato, P. Scarfato, L. Scatteia and D. Acierno, Polymer, 2004, 45, 3487.

Significant Assessment of Nanocomposites’ Combustion Behaviour

159

18. T.D. Fornes, P.J. Yoon, H. Keskkula and D.R. Paul, Polymer, 2001, 42, 9929. 19. V. Babrauskas, Ignition Handbook, Fire Science Publishers, Issaquah, 2003. 20. M. Zanetti, T. Kashiwagi, L. Falqui and G. Camino, Chem. Mater., 2002, 14, 881. 21. T. Kashiwagi and J.W. Gilman, in Fire retardancy of polymeric materials, ed. A. F. Grand, C. A. Wilkie, Marcel Dekker Inc., New York, 2000, p. 353.

CHAPTER 11

Phosphorus-Based Epoxy Resin–Nanoclay Composites JIANWEI HAO, YANBING XIONG AND NA WU National Laboratory of Flame Retardant Materials, School of Materials Science and Technology, Beijing Institute of Technology, Beijing, 100081, People’s Republic of China

11.1 Introduction Non-halogen flame retardant epoxy resins are in great demand for manufacturing printed circuit boards in electrical and electronic (E&E) equipment because recently the use of halogen flame retardants in plastic materials has raised many questions about their impact on the health and safety of both and the environment. In recent years significant attention has focused on the research on active phosphorus-containing flame retardants, nanofiller and their synergistic effect for epoxy resins.1–4 This especially worthy research has two aspects. Firstly, reactive-type phosphorus-containing flame retardants are of interest because they are able to sustain the flame retardance for a longer period, and reduce the influence on physical, mechanical and artifactual properties of polymers. Secondly, the synergistic effect between nanoclays and flame retardants can simultaneously improve the combustion performance and physical properties of the resin. Jeng et al. reported that synthesizing an epoxy resin that contained resorcinol diphosphate with 6.19 weight percent (wt%) of phosphorus resulted in epoxy resins that exhibited a high limiting oxygen index (LOI) value of 32% and high char yield of nearly 30% at 850 1C under nitrogen.5 Mauerer prepared epoxy Fire Retardancy of Polymers: New Strategies and Mechanisms Edited by T Richard Hull and Baljinder K Kandola r Royal Society of Chemistry 2009 Published by the Royal Society of Chemistry, www.rsc.org

160

Phosphorus-Based Epoxy Resin–Nanoclay Composites

161

resins that contained aminophenyl phosphate by reacting tris-(3-aminophenyl) phosphate (TAPP), a cross-linker and flame retardant, with curing agents like dicyandiamide or 4,4-diaminodiphenylsulfone.6 Laminates made of this aminophenyl phosphate cross-linked epoxy resin achieved the UL-94 V-1 classification, and also met the standard requirements for water absorption, heat stability and high glass-transition temperature. Hsiue et al. reported that curing a phosphorus-containing epoxy with an aminopropylpolydimethylsiloxane (PDMS) resulted in epoxy resins that have a high LOI value of 45%, hardly achievable using solely phosphorous compounds.7 Wu et al. researched the synergistic effects of phosphorus–silicon and nitrogen–silicon in the epoxy resins of silicon-containing epoxy compounds or nanoscale colloidal silica, cured with phosphorus- and melamine-containing agents.8 Addition of silicon compounds significantly enhanced the thermal stability, LOI values and char yields of the cured epoxy resins. Levchik and Weil1,2 have given a detailed overview of the recent literature on reactive flame retardants in epoxy resins, including on 9,10-dihydro-9-oxa10-phosphaphenanthrene-10-oxide (DOPO) compounds. DOPO was reacted with quinine, naphthalene and itaconic acid, respectively, and the resultant DOPO-containing compounds or condensates could be used as reactive flame retardants for bisphenol A epoxy resin or multifunctional novolac epoxy resin.1,2,9 The rigid cyclic side-chain structure of DOPO gave the phosphoruscontaining epoxy a higher glass-transition temperature and flexural modulus, and better thermal stability than the regular tetrabromobisphenol A epoxy resin. In the work reported in this chapter, the reactive-type organophosphorus flame retardant agent, DOPO was reacted with bisphenol A epoxy resin to form a phosphorus-based epoxide (EPO-P). This (EPO-P) was mixed with and without an organophilic clay (Cloisite 30A) in organic solvent to prepare phosphorus-containing epoxy–nanoclay (EP–P–nano) composites and a phosphorus-containing epoxy resin (EP–P), respectively. It was used in combination with a 4,4 0 -diaminodiphenyl-methan (DDM, a curing agent). The evidence for EP–P–nano composite formation was examined by X-ray diffraction (XRD) and transmission electron microscopy (TEM). The flame retardant properties of the EP–P–nano composites were measured by LOI and a UL-94 vertical flame test. The tensile and impact strengths of the composites were tested according to the standard methods.

11.2 Experimental 11.2.1

Materials

Bisphenol A epoxide (E-51, epoxy equivalent weight ¼ 209 g mol
1) was obtained from Wuxi Resins Factory in China. DOPO was obtained from Kunshan Zhonghong Chemical Inc. Curing agent DDM was supplied by Tianjin Guangfu Institute of Fine Chemical Industry. Cloisite 30A, an organophilic clay, was supplied by Southern Clay Products Inc.

162

Chapter 11 O

CH2

CH3 CH CH2 O

CH3

O O CH2 CH

C

XO

CH2 +

CH3

O

P

C

OX

CH3

O

H

Epoxide

O X=

EPO-P

DOPO

CH2 CH

OH CH2

or

CH2 CH CH2

P

O

O

Scheme 11.1

11.2.2

The reaction of epoxide and DOPO.

Preparation of Phosphorus-Based Epoxide

E-51 and DOPO in predetermined stoichiometry ratios were mixed and stirred to form a homogeneous solution at 120 1C for 1 hour, then reacted at 160 1C for 5 hours, according to Scheme 11.1 to obtain the transparent yellow EPO–P that contained 1.0 wt% of phosphorus. The epoxy equivalent weight of EPO–P is 241.6 g mol
1, which was measured by a chemical titration method.10

11.2.3

Preparation of Phosphorus-Based Epoxy–Nanoclay Composites

The EPO–P (phosphorus content 1.0 wt%) was mixed with organophilic clay dispersed in N,N 0 -dimethylformamide, at 80 1C for 1 hour, then DDM was added and stirred at 80 1C for 2 hours to obtain the cured EP–P–nano composites. Samples of EP–P–nano composites that contained 1, 3 and 5 wt% of organophilic clay were cured at 80 1C for 2 hours and post-cured at 120 1C for 12 hours.

11.2.4

Characterization

Fourier transform infrared (FTIR) spectra of a phosphorus-containing epoxide liquid film sample were obtained with a BRUKER Vector 22 spectrometer. XRD analyses were performed using a RINT-2400 diffractometer with Cu radiation (50 kV, 40 mA). The scanning speed and step size were 0.11 min
1 and 0.11, respectively. TEM images of EP–P–nano composites were taken with a HITACHI-H-800 at an acceleration voltage of 200 kV. The LOIs were measured using a FTA instrument according to ASTM D 2863-91. The vertical burn (VB) UL-94 test was conducted using a CZF-2 type instrument according to ANSI/UL 94-2006. The tensile strengths were measured following the GB/T 1040.2-2006 standard by a DXLL-5000 test apparatus and the impact strengths following the GB/T 1843-1996 standard by a XJ-502 impact test apparatus.

163

Phosphorus-Based Epoxy Resin–Nanoclay Composites

11.3 Results and Discussion 11.3.1

Structure of EPO–P Analysis with FTIR

The infrared spectra of DOPO and the EPO–P shown in Figures 11.1 and 11.2, respectively, were analyzed to prove that the phosphorus was introduced into the molecule of epoxide. The characteristic bands corresponding to DOPO and the EPO–P are given in Table 11.1. DOPO gave adsorption at 2438 cm
1, which corresponds to P–H stretching vibrations. Figure 11.2 shows that the peak of the P–H characteristic bands disappeared in EPO–P, which indicates that the phosphorus has been bonded on the epoxide molecule.

11.3.2

Structure of EP–P–nano Composites Analysis with XRD and TEM

XRD provides information on the changes in the inter-layer spacing of the clay upon the formation of a nanocomposite. XRD results of the organophilic clay and the EP–P–nano composites with 1 wt% of phosphorus and 3 wt% of Cloisite 30A clay cured at different temperatures are shown in Figure 11.3. It can be seen that for Cloisite 30A clay, 2y ¼ 4.71 and the d-spacing is 18.8 A˚. For EP–P–nano composite samples, the peak 001 has shifted to a lower value of 2y, the d-spacing is 36.0 A˚ and the peak is broadened. This indicates that the 100

Transmittance/%

80

60

3064 =CH(Ar)

2438 P-H

1604 C=C(Ar)

40

1481 P-Ph

1208 906 P-O-Ph

20 1244 P=O 0 4000

3500

3000

2500

2000

1500

Wavenumber/cm-1

Figure 11.1

FTIR spectra of DOPO.

1000

500

0

164

Chapter 11 100

Transmittance/%

80

3054 =CH(Ar)

60

764 P-C

1609 C=C(Ar)

40

1461 P-Ph 20

1186 910.25 P-O-Ph

2873 -CH2

1241 P=O

2975 0 4000

3500

3000

2500

2000

1500

1000

500

0

Wavenumber/cm-1

Figure 11.2 FTIR spectra of EPO–P.

Table 11.1

The characteristic FTIR bands of DOPO and EPO–P. The characteristic bands (cm
1)

Specimens

P–H

P–Ph

P¼O

P–O–Ph

DOPO EPO–P

2438 Disappear

1481 1461

1244 1241

1208, 906 1188, 910

EP–P–nano composite has a mixture of intercalated and exfoliated structures, and the curing temperature has no influence on the formation of nanocomposites. TEM provides an actual image of the clay layers and helps to identify morphology of the nanocomposite. Figure 11.4(a) shows the TEM images for the mixture of EP–P and 3 wt% Cloisite 30A˚ clay (EP–P/clay), where no intercalated clay layers can be seen. The TEM image of the EP–P–nano composite (3 wt% clay) in Figure 11.4(b) indicates that a mixed intercalated and exfoliated structure has been formed for this sample.

11.3.3

Combustion Performance and Mechanical Properties

The LOI determines the flame-extinguishing properties of the polymer materials under controlled test conditions. The VB test, UL-94, determines the

165

Intensity

Phosphorus-Based Epoxy Resin–Nanoclay Composites

a - Cloisite 30 A b - 120°C cured nano - composite c - 150°C cured nano - composite d - 80°C cured nano - composite 0

2

4

6

8

10

2θ (deg)

Figure 11.3 XRD patterns of EP–P–nano composites at different curing temperature.

a

b

100 nm

Figure 11.4

100 nm

TEM images for the mixture of a) EP–P/clay and b) EP–P–nano composite.

upward burning properties of the polymer materials. LOI values and UL-94 testing results of the epoxy-clay nanocomposite without phosphorus (EP– nano) and with phosphorus (EP–P–nano) are given in Table 11.2. The phosphorus content was fixed at 1 wt% in the EP– and EP–P–nano specimens. For sample EP–nano, containing only 3 wt% clay alone and no phosphorus, the LOI value was found to be 27.8%, which is a little higher than the 25.2% of the virgin epoxy resin (EP) specimen. When clay (o5 wt%) was added into EP–P

166

Table 11.2

Chapter 11

Data of combustion performance of epoxy resin specimens.

Specimens

P (wt%)

Clay (wt%)

LOI (%)

UL-94 (3 mm)

EP EP–P EP–nano

– 1 – 1 1 1 1 1

– – 3 1 2 3 4 5

25.2 29.8 27.8 30.1 31.1 32.1 30.4 28.6

V-2 V-1 V-2 V-1 V-0 V-0 V-1 V-1

EP–P–nano

Table 11.3

Mechanical properties of epoxy resin specimens.

EP EP–P EP–P–nano, 3 wt% of clay

Tensile strength (MPa)

Elongation at break (%)

Impact strength (kJ m
2)

60.2 69.7 72.3

3.6 4.2 4.5

13.3 16.8 25.4

with 1 wt% phosphorus, the LOI values and UL-94 ratings of EP–P–nano composites increased significantly. It can be seen that the LOI and UL-94 ratings of the EP–P–nano composite that contained 3 wt% of clay are 32.1% and V-0, respectively, which are the best flame retardant properties obtained. These results show that clay or phosphorus alone cannot effectively enhance the flame retardant properties of EP. In other words, a synergistic effect in flame retardancy can be seen when both phosphorus and clay are present in EP–P– nano composite samples, which indicates synergism between phosphorus and nanoclay. This synergism can be explained on the basis that phosphorus helps in the char formation and enrichment during combustion of the resin and the nanoclays enhance the gas barrier properties of the char, hence both acting together to exhibit the synergistic effect. Notice also that when the loading level of the clay increases to 4 wt% and 5 wt%, the LOI values and UL-94 rating of EP–P–nano composites decrease. This decrease in flame retardancy could be explained because when the EP–P– nano composites contain more than 4% clay, the cross-linking density of the EP–P–nanoclay composites during curing is reduced. The other reason could be that the clay at higher concentrations is not well dispersed in the composites, which results in a reduction in the gas-barrier properties of the char layer during burning. Table 11.3 presents the data of mechanical properties for the samples of EP, EP–P and EP–P–nano composite with 3 wt.% clay. It can be seen that the tensile strength and elongation of EP–P–nano composite increased from 69.7 MPa and 4.2% in EP to 72.3 MPa and 4.5%, respectively. Especially, the

Phosphorus-Based Epoxy Resin–Nanoclay Composites

167

impact strength of EP–P–nano composite increases significantly, i.e., from 13.3 kJ m
2 in EP to 25.4 kJ m
2. It is the formation of the intercalated and exfoliated structure in the EP–P–nano composite that contains 3 wt.% clay which leads to reinforcement of the epoxy substrate. The results indicate that flame retardant, phosphorus-based epoxy resin–nanocomposites could have increased tensile and impact strengths.

11.4 Conclusion DOPO was successfully introduced into an epoxy compound with 1 wt% phosphorus. The phosphorus-containing epoxy–nanoclay composites with various organophilic clay content were prepared. The results showed that the phosphorus-based flame retardant combined with nanocomposite is an effective way to improve flame retardancy and physical properties of the epoxy resin. The LOI and UL-94 ratings of the EP–P–nano composite with 3 wt% clay were 32.1% and V-0 rating, respectively. The Notched Izod impact strength of EP–P–nano reached 25.4 kJ m
2, increasing by 50% compared to the EP–P sample.

References 1. S. Levchik, A. Piotrowski, E. Weil and Q. Yao, Polym. Degrad. Stab., 2005, 88, 57. 2. S. Levchik and E. Weil, Polym. Int., 2004, 53, 1901. 3. A. Morgan, Polym. Adv. Technol., 2006, 17, 206. 4. Y. Xiong, J. Hao, Proceeding of the International Polymer Materials Engineering Conference 2005, Shanghai, China, September 18–21, 2005. 5. R. Jeng, S. Shau and J. Lin, Eur. Polym. J., 2002, 38, 683. 6. O. Mauerer, Polym. Degrad. Stab., 2005, 88, 70. 7. G. Hsiue and Y. Liu et al., J. Appl. Polym. Sci., 2000, 78, 1. 8. C. Wu, Y. Liu and Y. Chiu, Polymer, 2002, 43, 4277. 9. C. Wang and M. Lee, Polymer, 2000, 41, 3631. 10. J. Hao, Y. Xiong and T. Zhang, Transactions of the Beijing Institute of Technology, 2006, 26, 279.

CHAPTER 12

Study of the Relationship Between Flammability and Melt Rheological Properties of Flame-Retarded Poly(Butylene Terephthalate) Containing Nanoclays S. NAZARE,a T. R. HULL,b B. BISWAS,a F. SAMYN,c S. BOURBIGOT,c C. JAMA,c A. CASTROVINCI,d A. FINAd AND G. CAMINOd a

Centre for Materials Research and Innovation, University of Bolton, Deane Campus, Bolton, BL3 5AB, UK; b Centre for Fire Hazards Science, University of Central Lancashire, Preston, PR1 2HE, UK; c Proce´de´s d’Elaboration de Reveˆtements Fonctionnels (PERF), LSPES, UMR-CNRS 8008, ENSCL, BP 90108, 59650 Villeneuve d’Ascq, France; d Politecnico di Torino Sede di Alessandria-Centro di Cultura per l’Ingegneria delle Materie Plastiche, Viale Teresa Michel 5 15100, Alessandria, Italy

12.1 Introduction Recent studies on a new class of flame retardant (FR) systems that contain nanoclay and conventional FR microparticles have shown that the threshold concentration of FR required to achieve acceptable levels of flame retardancy Fire Retardancy of Polymers: New Strategies and Mechanisms Edited by T Richard Hull and Baljinder K Kandola r Royal Society of Chemistry 2009 Published by the Royal Society of Chemistry, www.rsc.org

168

169

Study of the Relationship Between Flammability and Melt Rheological Properties 1

can be significantly reduced in the presence of nanoclay. Bourbigot et al. have observed synergistic effects while incorporating nanofillers into intumescent formulations. They proposed that the reactivity of nanofillers with the intumescent FR modifies the physical behaviour of intumescent char during burning. In multicomponent polymer formulations that contain FR microparticles and inorganic nanoparticles, research has shown that the structure of the interphase (IP) strongly affects the flame retardancy and mechanical properties of the polymer system.2 The formation and structure of the IP is, however, governed by the interaction between solid–solid and solid–liquid phases. Nanoclays with different structural morphologies and organic surface treatments could interact differently with the FR microparticles, and thus result in materials with distinct physical properties. The structural morphology of the dispersed phase in the polymer strongly affects the rheological properties of the polymer system, which can sequentially alter the burning behaviour of the polymer composite.3 Therefore two different nanoclays with different structures were chosen. Cloisite 30B (CL30B) is a montmorillonite clay modified with a quaternary ammonium salt, which has a layered structure consisting of two tetrahedral silicate sheets that sandwich a central octahedral sheet. The aspect ratio of montmorillonite is very high, with a specific surface area of 750 m2 g
1. Sepiolite (SP), also a member of the same 2:1 phyllosilicate group, is a non-swelling clay with needlelike morphology. Chemically, SP is a microcrystalline-hydrated magnesium silicate with the unit cell formula of Si12O30Mg8(OH)4.8H2O.4 The SP structure consists of a magnesium octahedral sheet between two layers of silica tetrahedrons, which extend as a continuous layer with an inversion of the apical ends every six units. This inversion results in the formation of a discontinuous octahedral sheet, which allows the formation of rectangular tunnels growing in the direction of needle axis.5 The nanostructured tunnels measure approximately 0.35  1.06 nm2 in cross section and are filled with zeolitic water. The specific surface area of SP is 300 m2 g
1  10 m2 g
1 and the contact areas between the needles are both smaller than the specific surface area and contact area between the clay platelets of montmorillonite. The lower contact area between the needles facilitates dispersion of SP. Commercially, poly(butylene terephthalate) (PBT) is often rendered FR using halogen-containing additives and a synergist. However, environmental issues mean that halogenated systems are fast being replaced by additive or reactive flame retardant systems. Different flame retardant systems for PBT and thermal decomposition and combustion mechanisms of flame retarded PBT were recently reviewed by Levchik and Weil.6,7 In the present work, interactions between flame retardant microparticles and inorganic nanoparticles (of different morphologies) dispersed in PBT are examined using rheology, and changes in crystallinity and hence melting behaviour are studied using differential scanning calorimetry. Thermo-analytical studies were carried out to examine the effect of changed rheology on thermal decomposition of the polymer composites. Viscosity measurements as a function of temperature were carried out to obtain information about the interactions of the components and

170

Chapter 12

the processes that take place during heating. Finally, cone calorimetric experiments were performed to study the effects of changed melt rheological behaviour on the fire behaviour of PBT formulations.

12.2 Experimental 12.2.1

Materials

 Polymer – PBT, Celanex 2000-2 Natur, supplied by Ticona;  Nanofiller 1 – CL30B, natural montmorillonite modified with methyl, tallow, bis-2-hydroxyethyl, quaternary ammonium chloride (MT2EtOH), supplied by Southern Clay Products, USA;  Nanofiller 2 – SP amine, surface modified with benzyl methyl dihydrogenated tallow ammonium salt, supplied by Tolsa, Spain;  FR – phosphinate salt, Exolit OP1240, supplied by Clariant, Germany.

12.2.2

Sample Preparation

Compounds were prepared by melt blending in a Leistritz ZSE 27 co-rotating intermeshing twin-screw extruder. Screw speed was set to 200 revolution per minute (rpm) and mass flux at 10 kg h
1. Screw profile and temperature profile used for compounding PBT materials are shown in Figure 12.1. The polymer was loaded in the main feed and filler added to the molten polymer by means of a gravimetric side-feeder. The extruded materials are cooled in water and then pelletized. Samples in the form of powder, films and slabs were prepared for appropriate testing. Sample description and mass percentages of various components in the formulations are given in Table 12.1.

Figure 12.1

Screw and temperature profiles for processing PBT materials.

Study of the Relationship Between Flammability and Melt Rheological Properties

Table 12.1

171

Mass percentages of various components in the formulations.

Descriptive codes

Resin (%)

FR (%)

Nanoclay (%)

PBT PBT–CL30B PBT–SP PBT–FR PBT–FR–CL 30B PBT–FR–SP

100 95 95 82 77 77

– – – 18 18 18

– 5 5 – 5 5

12.2.3

Characterization and Testing

Conventionally, nanocomposite structure(s) in a polymer matrix can be identified by monitoring the position, shape and the intensity of the basal spacing in the lower 2y region (2–101) of the X-ray diffractogram (XRD). However, the peak in the XRD of SP originates from the d-spacing between the SP tunnels and not from the separation between the needles.4 Therefore, XRD is not a suitable technique to characterize the dispersion of SP in the polymer matrix. Recently, a rheological method was developed to characterize the nanodispersion of all kinds of plate-like, fibrous or dendritic filler materials with high aspect ratios.8 This was used to characterize the nanostructures of the samples in the current study. A Polymer Laboratories DSC was used to determine the influence of the morphological structure of the nanofillers on the crystallization behaviour of PBT. The crystallinity (Xc %) for all the samples was calculated such that:

X c% ¼

DH m  100 DH f

where DHm is the enthalpy of melting and DHf is enthalpy of fusion. The theoretical value of DHf for a 100% crystalline PBT has been taken as 140 J g
1.9 Simultaneous differential thermal analysis (DTA) and thermogravimetric analysis (TGA) was performed using an SDT 2960 TA instrument under flowing air (50 ml min
1) and at a heating rate of 10 K min
1 on 10 mg sample masses. Rheological measurements were carried out on 1 mm thick samples at 240 1C using a Dynamic Analyser Rheometer RDA II from Rheometrics. A parallel plate geometry with plate diameter 25 mm was used to conduct dynamic frequency sweep experiments. Furthermore, the changes in melt rheological behaviour of polymer composites over a temperature range close to, and above, the degradation temperature were studied in a nitrogen atmosphere. The samples were heated from 300 to 530 1C with a heating rate of 15 1C min
1. The frequency of oscillation was kept constant at 10 rad s
1 and the strain amplitude at 10%. The burning behaviour of PBT formulations was studied using cone calorimetry (Fire Testing Technology Ltd, UK). 100  100  6 mm samples were exposed to an incident heat flux of 50 kW m
2 under ambient atmosphere.

172

Chapter 12

Table 12.2

Rheological properties of PBT formulations at 0.1 rad s
1.

Samples

|Z| (Pa)

G 0 (dyn cm
2)

G00 (dyn cm
2)

Type of composite

PBT PBT–Cl30B

7.4  103 8.0  105

6.5  100 9.4  104

7.4  102 9.4  104

PBT–SP PBT–FR PBT–FR–CL30B

7.7  102 6.9  103 6.2  105

3.0  101 3.9  101 7.4  104

3.0  101 3.9  101 4.1  104

PBT–FR–SP

2.7  105

3.1  104

3.1  104

– Intercalated nanocomposite Microcomposite Microcomposite Intercalated nanocomposite Intercalated nanocomposite

12.3 Results and Discussion 12.3.1

Nanodispersion

Viscosity curves for PBT polymer and its composites are shown in Figure 12.2(a), and a summary of the rheological properties in the low frequency region (at 0.1 rad s
1) for all the formulations studied are given in Table 12.2. In Figure 12.2(a) PBT shows perfect Newtonian behaviour over all the frequency range measured, giving a shear-thinning component of Z ¼ 0.02. Addition of 5% of CL30B to the polymer matrix shows a shift to non-Newtonian behaviour in the low-frequency region and pronounced shear thinning (Z ¼ 0.67) at higher frequencies. A significant increase in the complex viscosity at lower frequencies and pronounced shear thinning in the higher frequency region at low loading levels of 5% w/w is a characteristic feature of intercalated–exfoliated nanocomposite structures.10 Characterization of PBT–CL30B as an intercalated nanocomposite based on its rheological behaviour is in agreement with the XRD results,11,12 in which the characteristic peak of CL30B at 2y ¼ 4.51 corresponding to a d-spacing of 1.88 nm has moved to a lower value of 2y ¼ 2.21, which indicates a d-spacing of 4.0 nm. XRD analyses, confirmed by TEM,11,12 show that, although the d-spacing has increased, the CL30B has still maintained its ordered platelet structure to form an intercalated nanocomposite. In contrast, PBT–SP samples that contain 5% w/w of needle-like nanofiller (SP) do not show any change in rheological properties. The viscosity curve in Figure 12.2(a) for PBT–SP shows perfect Newtonian behaviour similar to that of the pure PBT sample, which suggests that SP remains in tactoid form or does not form a percolated superstructure of well-dispersed nanoparticles. Lack of confinement of polymer chains by one-dimensional needle-like SP particles in PBT–SP samples accounts for the perfect Newtonian behaviour similar to that of pure PBT. Owing to the weak interaction between the SP particles and the PBT polymer, the tethering of polymer chains by SP is not strong enough. Moreover, the change in the yield behaviour of polymer–clay nanocomposite in molten form depends largely on the surface area of the particulates. Note that the specific surface area of montmorillonite clay is 750 m2 g
1, whereas that of SP can be less than 300 m2 g
1. The higher specific surface area of CL30B

Study of the Relationship Between Flammability and Melt Rheological Properties

173

provides greater resistance to polymer chains and hence higher viscosity, especially at lower frequencies. The montmorillonite-based CL30B forms a classic ‘‘card-house’’ structure. The polymer-layered nanocomposite structure is instrumental in imparting solid-like viscoelastic properties to PBT–CL30B samples. The rheological properties of PBT–FR in Table 12.2 do not show substantial change with respect to those of pure PBT, despite 18% w/w loading of FR. This suggests that addition of microparticles up to 18% w/w does not affect the chain movement and hence the rheological behaviour of the polymer system, whereas 5% w/w of nano-dispersed clay particles significantly affects rheological properties of the polymer nanocomposite. However, addition of FR to the PBT–SP formulation resulted in a sizeable increase in the viscosity of the PBT– FR–SP sample and a noticeable increase in the shear thinning at higher frequencies, which suggests that the FR assists in increasing compatibility between polymer chains and SP needles. SP has a very high concentration of surface silanols spaced every 0.5 nm along the length of the needles which facilitate coupling reactions with polymer, organic surfactant and/or the flame retardant. This could probably lead to diffusion of small molecules within the SP needles, and thereby assist uniform dispersion of SPs within the polymer matrix. Solidlike or pseudo solid-like viscoelastic behaviour of the PBT–FR–SP formulation, as seen in Figure 12.2 and Table 12.2, can be attributed to enhanced dispersion of SP in the presence of FR. Viscosity values for PBT–CL30B and PBT–FR–Cl30B over the whole frequency range tested are comparable (see Table 12.2), which suggests that the confined structure of CL30B within the polymer matrix and the chain stiffness of PBT limits further widening of interlayer space in the presence of FR. Furthermore, hydroxyl groups in the CL30B interlayer have two effects on PBT that contains carboxyl groups. First, it favours intercalation of PBT chains and the formation of an intercalated nanocomposite structure. Second, the enhanced interaction of ammonium cation with the silicate surface is less favourable for replacement of the surface

1000000

1000000 100000

100000

(b)

10000 G″, Pa

Iη*I,Pa.s

(a)

10000

1000 100 10

1000 0

1

10

100

1 100

Key

PBT

Figure 12.2

PBT+CL 30B

1000

10000

100000 1000000

G″, Pa

Frequency, rad/s PBT+SP

PBT+FR

PBT+FR+CL30B

PBT+FR+SP

(a) Viscosity versus frequency and (b) G 0 versus G00 plots for all PBT formulations.

174

Chapter 12

contacts by PBT chains, which thereby limits extensive intercalation and further exfoliation of CL30B in the PBT matrix.13 The shear-thinning behaviour of both the samples that contain CL30B is very similar [see Figure 12.2(a)] with a shear-thinning component of Z ¼ 0.67 for PBT–CL30B and Z ¼ 0.64 for PBT– FR–CL30B. Furthermore, the so called Cole–Cole plots (log G 0 versus log G00 ) in Figure 12.2(b) may be used to further elucidate the morphological state of such multiphase polymer systems. Note from Figure 12.2(b) that inclusion of CL30B in the pristine polymer shows a profound influence on the log G 0 versus log G00 plots and hence the morphological state as compared to the pure polymer and flame-retarded polymer. Addition of flame retardant to the PBT–SP formulations also shows an upward shift in log G00 versus log G00 plots, which suggests a change in the morphological state of the polymer system. Nanodispersion gives rise to a notable increase in the degree of heterogeneity of the polymeric system, thereby decreasing the slope of log G 0 versus log G00 plots, compared to PBT, PBT–SP and PBT–FR samples. That the log G 0 versus log G00 plots in Figure 12.2(b) differ for different samples suggests these polymer systems can be regarded as different materials from a rheological point of view. The frequency-dependent behaviour of storage and loss moduli of a polymer system is also related to its morphological state in molten form. The storage and loss moduli curves plotted as a function of frequency for PBT and its composites are shown in Figure 12.3. The frequency dependence of storage and loss moduli of PBT, PBT–FR and PBT–SP shown in Figures 12.3(a), 12.3(c) and 12.3(e) suggests that the viscoelastic behaviour of pure polymer is dominated by viscous liquid behaviour (with G 0 o G00 over all the frequency range measured and no cross-over frequency). However, for the sample PBT–Cl30B [see Figure 12.3(b)], G 0 4 G00 in the lower frequency region suggests solid-like behaviour due to physical jamming of clay platelets. The cross-over frequency is noted at 19.9 rad s
1, after which the polymer system exhibits viscous liquid behaviour. Addition of FR reduces the cross-over frequency to 6.3 rad s
1 for the PBT–FR–CL30B formulation. For PBT–FR–SP formulations, the crossover frequency is noted at the lower frequency of 3.2 rad s
1, which indicates that the interaction between needle-like particles of SP and polymer chains is lost at lower shear rates, leading to relaxation of the polymer chains and hence viscous liquid behaviour.

12.3.2

Differential Scanning Calorimetry and Thermal Analysis

Calorimetric data for pure PBT and normalized (for actual polymer content) calorimetric values for all PBT formulations are given in Table 12.3. Note that the melting temperatures have remained unchanged. However, the samples that contain SP exhibit higher crystallization temperatures (Tc) than those of pure PBT and flame-retarded PBT, both with and without CL30B. The increased temperature of crystallization for PBT–SP may result from the reduced confinement effect of the one-dimensional needle-like SP clay particles, compared to that of the two-dimensional MMT platelets.14 Furthermore, the crystallization

175

Study of the Relationship Between Flammability and Melt Rheological Properties 1.E+07 1.E+05 1.E+04 1.E+03 1.E+02

PBT G' PBT G"

1.E+01

1 10 Frequency. rad/s

1.E+03 1.E+02

PBT+CL 30B G' PBT+CL 30B G" 0.1

100

1 10 Frequency. rad/s

100

1.E+07

1.E+07 (c)

1.E+06

(d)

1.E+06

1.E+05

G' and G" Pa

G' and G" Pa

wc = 19.9

1.E+04

1.E+00 0.1

1.E+04 1.E+03 1.E+02

PBT+FR G' PBT+FR G"

1.E+01 0.1

1 10 Frequency. rad/s

1.E+05

wc = 6.3 rads/s

1.E+04 1.E+03 PBT+FR+CL 30B G' PBT+FR+CL 30B G"

1.E+02 1.E+01

1.E+00 100

1.E+00 0.1

1.E+07

1.E+07

(e)

1.E+05 1.E+04 1.E+03 1.E+02 PBT+SP G' PBT+SP G"

1.E+01

1 10 Frequency. rad/s

1.E+05 1.E+04

wc = 3.2

1.E+03 1.E+02 PBT+FR+SP G' PBT+FR+SP G"

1.E+01 1.E+00

1.E+00 0.1

Figure 12.3

Table 12.3

1 10 Frequency. rad/s

100

(f )

1.E+06 G' and G" Pa

G' and G" Pa

1.E+05

1.E+01

1.E+00

1.E+06

(b)

1.E+06 G' and G" Pa

G' and G" Pa

1.E+07

(a)

1.E+06

0.1

100

1

10 Frequency. rad/s

100

Storage modulus and loss modulus of PBT and its composites.

Calorimetric data for PBT formulations.

Samples

Tm (1C)

Tc (1C)

DHm (J g
1)

DHc (J g
1)

xc (%)

PBT PBT–Cl30B PBT–SP PBT–FR PBT–FR–CL30B PBT–FR–SP

225 224 224 225 224 225

195 193 198 190 193 203

41 48 47 43 48 56

57 69 62 59 57 65

29 35 34 31 34 40

process starts much earlier in SP-containing samples, but the enthalpy of crystallization, DHc, is smaller than that of the PBT–CL30B formulation, which suggests the formation of larger crystals with fewer nucleating sites15 in the PBT–SP sample. Nanodispersed clay platelets in PBT–CL30B provide more heterophase nuclei and a larger surface area to increase DHc, but the triggering of the crystallization is slightly delayed.16 Addition of flame-retardant microparticles reduces the temperature of crystallization of PBT–FR formulation (Tc ¼ 190 1C), compared to that of the pure polymer (Tc ¼ 195 1C). Moreover,

176

Chapter 12

(a)

Mass Loss, %

100 80 60 40 20 0 0

Figure 12.4

200 400 600 Temperature, °C

800

Temperature Difference, °C/min

inclusion of CL30B in the flame-retarded PBT slightly shifts Tc to a higher temperature, but the enthalpy of crystallization is still lower than that for the PBT–FR sample. On the contrary, addition of SP to the PBT–FR sample significantly increases Tc and DHc of the resulting PBT–FR–SP sample, which suggests early onset of crystallization in the presence of SP particles. An increase in enthalpy of crystallization may be explained by improved dispersion of SP particles in the presence of flame-retardant particles, and hence enhanced interaction between SP particles and polymer chains. The normalized values for enthalpy of melting recorded during the second heating cycles are higher for PBT formulations that contain nanofillers, which suggests that greater resistance to melting is offered by the nanofillers. Enthalpy of melting is highest for the PBT–FR–SP sample (56 J g
1), confirming that the SP is nanodispersed in the presence of FR. The percent crystallinity for PBT– FR–SP is the highest of all the samples. The increase in crystallinity can be attributed to the nanodispersed SP needles providing heterophase nuclei. One of the most important property enhancements expected from formation of a polymer nanocomposite is that of thermal stability, either in the initial stages or in the final carbonaceous residues. The degradation of pure PBT in the presence of air proceeds through a free-radical mechanism. The TGA and DTA curves for pure PBT, PBT that contains CL–30B and SP are shown in Figure 12.4. The presence of nanoclays has no impact on the thermal stability of PBT below 400 1C. The clay layers act as a mass-transport barrier to the volatile products generated during decomposition, which increases thermal stability. However, there are also catalytically active centres in the clay layers, such as those around hydroxyl groups, which might accelerate the decomposition of PBT.16 Above 400 1C both the clays improve the thermal stability of PBT and give rise to similar yields of carbonaceous char at high temperatures. Although the TGA curves for PBT–CL30B and PBT–SP show a similar trend, the DTA curves are quite different. The small exothermic peak at 279 1C for PBT–CL30B suggests decomposition of the organic modifier, whereas the organic modifier on SP is stable up to 300 1C. The DTA curve for PBT–SP shows an exothermic peak at 352 1C, which could be caused by degradation of the amine group,

12 10 8 6 4 2 0

PBT PBT+SP PBT+CL 30B

(b)

0

200 400 600 Temperature, °C

TGA and DTA responses in air for PBT-based materials.

800

177

Study of the Relationship Between Flammability and Melt Rheological Properties

followed by an endotherm that could be attributed to dehydration in which SP loses half of its coordinated water.5 The main exothermic peak for pure PBT at 417 1C, which represents the release of volatiles, is much smaller for PBT– CL30B and PBT–SP samples, which could probably result from the barrier effect of nanoclays. Inclusion of FR in the formulations that contain two different clays (not shown here) does not have any synergistic effect on the thermal stability of PBT.

12.3.3

Melt Viscosity

Viscosity versus temperature curves for PBT-based materials are given in Figure 12.5. The expanded scale within Figure 12.5 shows that the viscosity of neat PBT reduces to near to zero up to 435 1C, caused by melting and then complete decomposition of the polymer. A sharp increase in viscosity of PBT samples above 435 1C can be attributed to the presence of solid carbonaceous residue. Viscosity measurements beyond 435 1C for pure PBT were not possible for instrumental limitations. The viscous modulus of the PBT–SP formulation is greater by a factor of 10 than that of pure PBT. This increase in viscosity of the PBT–SP formulation over the temperature range 300–415 1C, despite a small (5% w/w) loading of SP, results from reinforcement of the polymer matrix by needle-like nanoparticles of SP. However, this effect of adding SP is not seen in the viscoelastic properties measured at 240 1C. This suggests that, at higher temperatures, dispersion of SP is improved, which results in an increased viscosity of PBT–SP. However, this increase in viscosity is not sufficient to prevent melt dripping of the sample when exposed to an external heat flux or flame. Above 420 1C the viscosity of PBT–SP falls to near zero, through degradation of the polymer. A sharp increase in the viscous modulus at 500 1C

Iη*I, Pa.s

25000

Iη*I, Pa.s

20000 15000

800 700 600 500 400 300 200 100 0 300

350

400

450

500

550

Temperature, °C

10000 5000 0 300

350

400

450

500

550

Temperature, °C

Figure 12.5

PBT

PBT+CL 30B

PBT+SP

PBT+FR

PBT+FR+CL 30B

PBT+FR+SP

Viscosity versus temperature curves for PBT formulations.

178

Chapter 12

could be attributed to the formation of a solid inorganic char. Note from the inset plot in Figure 12.5 that the degradation step of PBT–SP is delayed compared to those of both the pure and the flame-retarded PBT. As seen in Figure 12.5, the increased viscosity of the PBT–CL30B sample, compared to those of the PBT, PBT–SP and PBT–FR formulations, over the temperature range 300–350 1C suggests increased resistance to melt dripping. Figure 12.5 also shows that above 350 1C, the viscosity for the PBT–CL30B sample does not come close to zero until 425 1C, which suggests further resistance to melting over the temperature range 350–425 1C. Changes in viscosity with increasing temperature for PBT–FR up to 320 1C are similar to those in the pure PBT sample. However, at 335 1C, a viscosity peak appears which can be assigned to flame-retardant activity in the presence of the P-based intumescent flame retardant. This peak gradually levels to zero around 435 1C, which could be caused by the formation of phosphoric acid species from the thermal decomposition of the phosphinate. A sharp increase in viscosity and subsequent stability at higher temperatures for the PBT–FR formulation can be attributed to the enhanced formation of char in the presence of the FR. Finally, addition of 5% w/w of CL30B and SP to the PBT–FR formulation dramatically increases viscosity in the resulting PBT–FR–CL30B and PBT–FR– SP samples. A gradual decrease in viscosity values of the PBT–FR–CL30B and PBT–FR–SP formulations above 325 1C could result from the formation of phosphoric acid species, as mentioned above. Moreover, the polyphosphoric acid may react with the surfactant of the nanoclay and thereby collapse the nanostructure, and thus result in lower viscosities of PBT–FR–CL30B and PBT–FR– SP. However, the appearance of a shoulder at 360 1C (for the PBT–FR–SP formulation) and a viscosity peak at 415 1C (for the PBT–FR–CL30B formulation) suggests the formation of a porous carbonaceous char which subsequently collapses. This reduces viscosity to near zero for both PBT–FR–SP and PBT–FR– CL30B. A sharp increase in viscosity of PBT–FR–SP above 410 1C may be caused by the formation of a char that is reinforced with needle-like nanoparticles. Addition of CL30B to PBT–FR has a slightly different effect on the viscosity of the resultant formulation than does addition of SP. As seen in Figure 12.5, the reduction in viscosity is more gradual and prolonged, compared to that of the PBT–FR–SP sample. The barrier effect of nanodispersed clay platelets in the polymer matrix means the degradation step of the PBT–FR–CL30B formulation is delayed compared to that of the PBT–FR–SP formulation. The final charring process starts at 500 1C, as opposed to 410 1C, for the PBT–FR–SP sample. From the above discussion, it can be concluded that the PBT–FR–CL30B formulation might be expected to show the better fire performance because of the increased viscosity and thermal stability through the presence of nanoclay.

12.3.4

Flammability

The cone data obtained at 50 kW m
2 and given in Table 12.4 shows significant differences for various PBT formulations. Most importantly, and of more

Cone calorimetric results at 50 kW m–2 heat flux for all PBT formulations.a

Sample

PHRR TTI (s) (kW m
2)

AHRRb (kW m
2)

THRb (MJ m
2)

FIGRA (kW s
1)

Hcb (MJ kg
1)

Char residueb (%)

CO (g g
1)

CO2 (g g
1)

PBT PBT–Cl30B PBT–SP PBT–FR PBT–FR–CL30B PBT–FR–SP

64 51 44 42 37 40

229 177 191 140 110 116

138 106 115 85 66 70

1.9 2.1 3.0 1.4 1.3 1.9

22 19 21 15 13 16

27 34 37 32 41 49

0.26 0.14 0.13 0.17 0.17 0.21

2.38 1.85 2.16 1.78 0.84 1.19

a b

597 279 332 250 165 163

AHRR, average heat-release rate; FIGRA, fire growth rate; PHRR, peak heat-release rate; THR, total heat release. Values at 600 s.

Study of the Relationship Between Flammability and Melt Rheological Properties

Table 12.4

179

180

Chapter 12

significance to this work, is the time to ignition (TTI). A critical surface temperature for ignition is close to being accepted as a material property, and the time to reach this temperature (TTI) is a function of the heat transfers.17 TTI for the neat PBT is greater than the average of the nano- or FR-containing formulations. Several factors influence the ignition delay time. However, based on our rheological studies and observations, we propose a hypothesis that an increasing viscosity decreases thermal conductivity, essentially by flowing of the molten polymer, and thus results in accumulation of heat at the surface of the sample exposed to an incident heat flux. Furthermore, the thermal properties (krc) of the solid material are relatively easy to define and measure, but as the rheometric data show, most samples are somewhat molten at their ignition temperature. The increased surface temperature of the sample with higher viscosity means that this sample reaches ignition temperature more quickly than the sample with low viscosity. Based on this argument, the pure PBT would flow and bubble, and thus allow the whole sample to reach thermal equilibrium and thus increase the TTI. Once the bulk PBT reached the ignition temperature, the burning rate would be more rapid (see Figure 12.6) and lead to higher values of PHRR, FIGRA and AHRR, as seen in Table 12.4. For the samples that contain only nanofiller, the reduction in TTI can be ascribed to the increased viscosity at the ignition temperature, which results in a higher surface temperature (but a lower bulk temperature). Of the two nanofillers, the SP-containing samples show shorter TTIs than CL30B-containing formulations. This is in contrast to the above hypothesis, since the PBT–CL30B sample with higher viscosity shows an increased TTI compared to the PBT–SP sample with lower viscosity. The increased TTI in

700 600 PBT PBT+ CL 30B PBT+SP PBT+ FR PBT+FR+ CL 30B PBT+FR+SP

HRR, kW/m2

500 400 300 200 100 0 0

200

400

600

800

1000

1200

1400

Time, s

Figure 12.6

HRR as a function of time for PBT formulations.

1600

1800

2000

181

Study of the Relationship Between Flammability and Melt Rheological Properties

PBT–CL30B can be attributed to several other factors, including adsorption of volatile products on larger surface areas of clay particles and the barrier effect of the plate-like CL30B. The early ignition of PBT–SP could also be caused by catalytic degradation of the SP amine and/or less-efficient barrier properties of SP clay. With increased viscosity, HRR is decreased to give lower PHRR and lower AHRR. The higher viscosity in the presence of nanoclay may also inhibit the escape of volatile products from the burning polymer into the flaming zone, and so reduce the HRR. THR values reported at 600 s are also reduced because of slower burning of samples that contain nanofillers. Furthermore, to study the effect of changed rheological properties on the flammability of PBT composites, the relationships between the intrinsic viscosity measured at 300 1C and various cone parameters are plotted in Figure 12.7. However, to eliminate the additional effect of FR, only PBT, PBT–CL30B and PBT–SP are compared. Moreover, the presence of FR would further obscure the effect of changed viscosity on the burning behaviour of the PBT–FR– CL30B and PBT–FR–SP samples. Figure 12.7(a) suggests that the TTI is related to viscosity, but that other factors, such as nanoparticle morphology or the ability of nanoparticles to act as a barrier, must also be involved. The PHRR decreases as the viscosity increases, especially between PBT and PBT–SP, but (as discussed earlier) the higher PHRR of PBT may be the consequence of the higher overall temperature of bulk polymer, compared to the higher surface

(a)

PHRR, kW/m2

TTI, s

700 70 60 50 40 30 20 10 0

(b)

600 500 400 300 200 100 0

10

100

1000

10000 100000 1000000

10

100

1000

23

40 35 30 25 20 15 10 5 0

(d)

(c) Hc, MJ/m2

Char residue, %

10000 100000 1000000

η, Pa.s

η, Pa.s 22 21 20 19 18 10

100

1000

10000 100000 1000000

10

100

1000

η, Pa.s PBT,

Figure 12.7

10000 100000 1000000

η, Pa.s PBT+CL 30B and

PBT+SP

Relationship between (a) TTI, (b) PHRR, (c) char residue and (d) Hc.

182

Chapter 12

temperature of the nanofilled PBT samples. Figure 12.7(c) shows that the char yield appears to be independent of viscosity, and is probably dependent on the processes that occur in the later stages of burning. That the char yield does not correlate with the THR suggests some inconsistencies in the burning behaviour. Furthermore, the modest decrease in heat of combustion with increase in viscosity implies a change in the gas-phase oxidation behaviour of the volatile products. Again, this is most likely to be a consequence of the cooler bulk of the nanofilled PBT materials, which results in incomplete gas-phase combustion and greater char formation. In summary, the plots in Figure 12.5 suggest that PBT formulations with higher viscosities exhibit improved post-ignition flameretardant properties.

12.4 Conclusions Rheological measurements suggest that one-dimensional, needle-like SP has a reduced confinement effect compared to the two-dimensional platelets of CL30B. This results in perfect Newtonian viscous behaviour of PBT–SP melts. This is also confirmed by calorimetric results, in which PBT–SP samples show higher crystallization temperatures and smaller enthalpies of crystallization compared to those of PBT–CL30B. This suggests the formation of larger crystals with fewer nucleating sites. The FR acts as a compatibilizer and facilitates better dispersion of SP to give higher melt viscosity for PBT–FR–SP formulations in the lower frequency region, and pronounced shear thinning at higher frequencies. The presence of FR in PBT–FR–CL30B formulations, however, does not affect their melt rheological properties. Despite bringing about changes in melt viscosity, melting and crystallization, the introduction of the clays, CL30B and SP, does not seem to alter the thermal degradation of PBT. In terms of melting behaviour, the viscosity measurements over a temperature ramp have shown that increased viscosity in the presence of nanoclay prevents dripping and flowing of polymer. In the cone calorimetric studies, this relates to shortening the TTI and a reduction in the rate of heat release. Furthermore, PBT formulations that contain CL30B show inhibited post-ignition combustion reactions, possibly because of physicochemical adsorption of volatile degradation products on the surface of silicates with higher specific surface area compared to those of their SP analogues.

Acknowledgements The authors gratefully acknowledge the financial support from the European Union through the Sixth Framework Programme Priority 3 NMP ‘‘PREDFIRE-NANO’’ (Contract No.: STREP 013998), and thank Dr Andy Prike from the University of Sheffield for his help with the rheology experiments.

Study of the Relationship Between Flammability and Melt Rheological Properties

183

References 1. S. Bourbigot, S. Duquesne, G. Fontaine, T. Turf, S. Bellayer, in Proceedings of The Eighteenth Annual BBC Conference, Stamford, Connecticut, May, 2007, 21–23. 2. S. Keszei, Sz. Matko´, Gy. Bertalan, P. Anna, Gy. Marosi and A. To´th, Eur. Polym. J., 2005, 41, 697. 3. J. Tung, R.K. Gupta, G.P. Simon, G.H. Edward and S.N. Bhattacharya, Polymer, 2005, 46, 10405. 4. A. Nohales, L. Solar, I. Porcar, C. Vallo and C.M. Go´mez, Eur. Polym. J., 2006, 42, 3093. 5. G. Tartaglione, D. Tabuani and G. Camino, Microporous Mesoporous Mater., 2008, 107(1–2), 161. 6. S.V. Levchik and E.D. Weil, Polymer International, 2004, 54, 11. 7. S.V. Levchik and E.D. Weil, A review on thermal decomposition and combustion of thermoplastic polyesters, 2004, 15(12), 691. 8. J. Zhao, A.B. Morgan and J.D. Harris, Rheological characterization of polystyrene–clay nanocomposites to compare the degree of exfoliation and dispersion, Polymer, 2005, 46, 8641. 9. G. Broza, Z. Kwiatkowska and R. K. Schulte, Processing and assessment of poly(butylene terepthalate) nanocomposites reinforced with oxidised single wall carbon nanotubes, Polymer, 2005, 46, 5860. 10. R. Wagener and T.J.G. Reisinger, Polymer, 2003, 44, 7513. 11. F. Samyn, S. Bourbigot, C. Jama, S. Bellayer, S. Nazare, T.R. Hull, A. Castrovinci, A. Fina and G. Camino, in preparation. 12. F. Samyn, S. Bourbigot, C. Jama, S. Bellayer, S. Nazare, T.R. Hull, A. Castrovinci, A. Fina and G. Camino, in preparation. 13. X. Li, T. Kang, W.J. Cho and C.S. Ha Lee, Macromol. Rapid Commun., 2001, 22, 1306. 14. S. Xie, S. Zhang, F. Wang, M. Yang, R. Se´gue´la and J.M. Lefebvre, Composites Science and Technology, 2007, 67, 2334. 15. X.Y. Tian, C.J. Ruan, P. Cui, W.T. Liu, J. Zheng, X. Zhang, X.Y. Yao, K. Zheng and Y. Li, Chem. Eng. Comm., 2007, 194, 205. 16. J. Xiao, Y. Hu, Z. Wang, Y. Tang, Z. Chen and W. Fan, Eur. Polym. J., 2005, 41, 1030. 17. B. Schartel and T.R. Hull, Fire and Materials, 31(5), 327.

CHAPTER 13

Thermal and Fire Performance of Flame-Retarded Epoxy Resin: Investigating Interaction Between Resorcinol Bis(Diphenyl Phosphate) and Epoxy Nanocomposites CHARALAMPOS KATSOULIS, EVERSON KANDARE AND BALJINDER K. KANDOLA Centre for Materials Research and Innovation, University of Bolton, BL3 5AB, UK

13.1 Introduction The flammability of polymeric materials has always been a great concern, because it limits their use in applications where fire safety is a key criterion. The use of both halogenated and non-halogenated conventional flame retardants has proved to be an effective solution to the problem, and resulted in materials with improved thermal stability and reduced flammability.1,2 However, studies have shown that the addition of high concentrations of flame retardants in a polymeric system often causes significant deterioration of the mechanical properties of the material. Moreover, environmental concerns have also been

Fire Retardancy of Polymers: New Strategies and Mechanisms Edited by T Richard Hull and Baljinder K Kandola r Royal Society of Chemistry 2009 Published by the Royal Society of Chemistry, www.rsc.org

184

Thermal and Fire Performance of Flame-Retarded Epoxy Resin

185

raised regarding the production of toxic gases during the combustion of polymers that contain halogenated flame retardants.1,3–5 In contrast, polymer nanocomposites have been attracting considerable attention as they offer improved thermal and fire properties not yet attained with conventional flame retardants.6–11 The most important characteristic of polymer–nanocomposites is that they can maintain their mechanical integrity while exposed to high heat-flux sources. These improvements can be achieved by using low additive contents [2–5 weight percent (wt%)] of nanoparticles such as layered silicate clays. The degree of dispersion of nanoparticles in the polymer matrix affects the enhancement in mechanical strength.12–16 Conventional flame retardants and nanoparticles when used individually are effective in improving some, but not all, of the fire-retardant and physical properties of polymers.17,18 The addition of at least two flame-retardant additives into a polymer matrix opens vast possible ways via which they can interact to result in an efficient flame-retardant system. When two or more additives are concomitantly interspersed into the polymer, their interactions may be additive, synergistic or antagonistic. Synergism means that the observed effect is greater than the additive, while an antagonistic effect is less than the additive.19 In this study the additive effect of resorcinol bis(diphenyl phosphate) (RDP) and nanoclay, Nanomer I.30E (I.30E), either individually or in combination, has been investigated. The phosphorus-containing additive RDP is an effective flame retardant; however, its low molecular weight renders it volatile, which causes partial loss from the polymer in the early stages of combustion. The reason for combining the phosphorus-containing additive RDP with the nanoclay is to have the flame retardant possibly intercalated between the nanolayers together with epoxy polymer chains so as to increase its residence time in the condensed phase during combustion. The long-term goal of this work is to develop flameretardant combinations that will be effective with respect to multiple flameretardant parameters at low concentrations.

13.2 Experimental 13.2.1

Materials

The epoxy resin, tetraglycidyl-4,4 0 -diaminodiphenylmethane (TGDDM), Araldite MY-721, and the curing agent, 4,4 0 -diaminodiphenyl sulfone (DDS), Aradur 9761, were supplied by Huntsman Corporation, Switzerland. RDP flame retardant was provided by Great Lakes, UK. The commercial grade organoclay I.30E was supplied by Nanocor Inc., while dichloromethane (CH2Cl2) was supplied by Aldrich-Sigma, UK. All the chemicals were used as received.

13.2.2

Sample Preparation and Characterization

Epoxy resin composites incorporating RDP and an organically modified clay, I.30E, were prepared by adding predetermined additive weight fractions to a

186

Chapter 13

Table 13.1

Mass percentages of various components in flame-retarded epoxy formulations. Mass (%)

Sample

Epoxy

RDP (P)a

I.30E

Epoxy Epoxy–RDP (95/5) Epoxy–RDP (90/10) Epoxy–RDP (85/15) Epoxy–RDP (80/20) Epoxy–RDP–I.30E (95/0/5) Epoxy–RDP–I.30E (90/5/5) Epoxy–RDP–I.30E (85/10/5) Epoxy–RDP–I.30E (80/15/5)

100 95 90 85 80 95 90 85 80

0(0.0) 5(0.5) 10(1.1) 15(1.6) 20(2.2) 0(0.0) 5(0.5) 10(1.1) 15(1.6)

0 0 0 0 0 5 5 5 5

a

Italicized values in parentheses represent the actual phosphorus mass fractional percent content. These values were calculated from the additive RDP weight fractions.

measured resin portion heated at 80 1C. A stoichiometric amount of the curing agent DDS (30% weight fraction with respect to epoxy formulation) and solvent CH2Cl2 (30.0 g, 0.4 mmol) were then slowly added with vigorous mechanical stirring for 30–40 minutes at room temperature to afford homogeneity. The resultant homogeneous mixture was degassed in a vacuum oven at 80 1C to evaporate the solvent and eliminate trapped air bubbles. Samples were cast in preformed aluminium dishes and cured at 180 1C for 3 hours. Details of constitutional percent loadings of respective additives are presented in Table 13.1. The morphology of the cured samples was investigated by using an X-ray diffractometer (Philips powder diffractometer) Cu Ka2 (l ¼ 1.54 A˚) radiation source at a scanning rate of 21 min
1 ranging from 21 to 601. Transmission electron microscopy (TEM) observations were also conducted using a Jeol JEM-2100 LaB6 microscope at an acceleration voltage of 200 kV with a current of approximately 100 mA.

13.2.3

Thermogravimetric Analysis

Thermogravimetric analysis (TGA) and differential thermal analysis (DTA) were performed on an SDT 2960 simultaneous DTA–TGA instrument from room temperature to 800 1C using 15  1 mg samples heated at 10 1C min
1 with air as the purge gas (flow rate, 100  5 ml min
1). The experiments were performed in triplicate and showed good reproducibility.

13.2.4

Flammability Tests

The flammability behaviours of flame-retarded resin laminates were determined via a UL-94 vertical test in accordance with protocol ASTM D 3801 with sample dimensions of 125 mm  12.5 mm  3 mm. The limiting oxygen index (LOI; Stanton-Redcroft Ltd) was determined in accordance with British

187

Thermal and Fire Performance of Flame-Retarded Epoxy Resin

standard method BS 2782; sample dimensions were 150 mm  12.5 mm  3 mm. A cone calorimeter (Fire Testing Technology Ltd, UK) was used to assess the time to ignition (TTI), heat release rate (HRR) as a function of time, total heat release (THR) and mass loss as a function of time. The test was carried out according to ISO 5660 standards, at an incident heat flux of 50 kW m
2. At least three round specimens of 56 mm diameter and 3 mm thickness were tested for each formulation.

13.3 Results and Discussion 13.3.1

XRD and TEM Analysis

XRD and TEM are techniques commonly used to identify the structural morphology of polymer nanocomposites, with observation of the extent of separation of clay layers and/or their exfoliation. The XRD patterns of the epoxy resin, epoxy–RDP–I.30E -(95/0/5) and epoxy–RDP–I.30E -(80/15/5) formulations are shown in Figure 13.1. The XRD patterns of epoxy formulations that contain 15% RDP and/or 5% I.30E by mass are similar to that of the unmodified resin. These patterns reveal the absence of basal reflections (00l ) at low angles (2–101), which suggests either complete delamination of the clay platelets (exfoliated structure) or very short-range order stacking of nanoclay platelets. The presence of the epoxy crystallographic peak centred at about 191 rules out low sensitivity of the instrument as the reason for the absence of 00l reflections from the nanoclay.

Intensity count (arb. units)

Intensity count (arb. units)

800

600

400

(c) 300

(b)

200 100

(a)

0 2

400

4

6 8 2 theta (°)

10

12

(c)

(b)

200

(a) 0 0

10

20

30

40

50

60

2 theta (°)

Figure 13.1

XRD data for (a) epoxy resin, (b) epoxy–RDP–I.30E (95/0/5) and (c) epoxy–RDP–I.30E (80/15/5) formulations. Insert shows magnified region between 2 and 101. Data are offset for clarity, but otherwise not scaled.

188

Chapter 13

(a)

(b)

tactoids

100 nm

50 nm

Figure 13.2

TEM images at low magnification (left) and high magnification (right) of an epoxy nanocomposite containing 5% I.30E clay together with 15% RDP.

TEM can be used to evaluate more directly whether some exfoliated phases exist within the polymer matrix. The low-magnification images provide information about the nanodispersion, while high-magnification images tell whether exfoliation and/or intercalation have been achieved. Figure 13.2(a) shows a TEM image of epoxy–RDP–I.30E -(80/15/5) at low magnification. The image reveals the presence of some regions that contain epoxy resin alone and dark circles that, perhaps, corresponds to clustered nanoclay particles (tactoids and agglomerates). The higher magnification image, Figure 13.2(b), reveals a nanoclay-occupied area that shows intercalated nanolayers with gallery spacings estimated at 5–10 nm. The overall picture shows that the modified nanoclay layers did not occupy the full volume with regions of epoxy visible. The absence of basal reflections at lower 2y values may result from a low degree of ordering of the nanoclay particles in the c-dimension.

13.3.2 13.3.2.1

Thermal Degradation Behaviour of Epoxy Resin and Its Composites Effect of RDP on Thermal Stability of Epoxy Resin

Figure 13.3(a) shows TG mass-loss curves for the unmodified epoxy resin and its RDP flame-retarded formulations. Thermal degradation of unmodified epoxy resin occurs primarily in three stages:20  the first a dehydration stage (100–240 1C) in which up to 3% mass loss is observed;  the second stage (240–470 1C) shows a 42% mass loss attributed to the depolymerization of polymeric chains to form a primary carbonaceous char;  the third stage is a char oxidation stage (470–680 1C) that contributes 55% mass loss and leaves no residual char at 800 1C.21,22

189

Thermal and Fire Performance of Flame-Retarded Epoxy Resin 100

Epoxy Epoxy/RDP-95/ 5 Epoxy/RDP-90/10 Epoxy/RDP-85/15 Epoxy/RDP-80/20

(a)

Mass %

80 60 40 20 0 (b)

∆ Mass %

20

Epoxy/RDP-95/ 5 Epoxy/RDP-90/10 Epoxy/RDP-85/15 Epoxy/RDP-80/20

10 0 -10 -20 50

150

250

350

450

550

650

750

Temperature (°C)

Figure 13.3 (a) TGA curves and (b) mass-difference curves between RDP flameretarded and unmodified epoxy resin versus temperature in flowing air.

Addition of RDP to the epoxy resin leads to a significant reduction in the onset degradation temperature of the resultant flame-retarded epoxy composites, measured as the temperature at which 10% mass loss occurs, T10 (Table 13.2). This suggests that the thermal stability of epoxy is compromised in the lower temperature regime after the addition of RDP. However, thermal stability for formulations that contain more than 5% RDP is improved from 425 1C upwards, as shown in Figure 13.3(a). When the epoxy resin is modified by the addition of 5% RDP, its thermal stability is only improved between 510 and 660 1C, which may be attributed to the low phosphorus content of RDP, calculated at 0.5 mass percentage of phosphorus (Table 13.1). Mass-difference curves (mass percent of epoxy–RDP samples minus mass percent of unmodified epoxy at the same degradation temperature) for epoxy and its flame-retarded composites are shown in Figure 13.3(b). The addition of 5% RDP destabilizes the epoxy resin at temperatures between 250–510 1C and beyond 660 1C, as shown by negative D mass percent values in those regions. A slight thermal stabilization is realized between 510 and 660 1C. Increasing the weight percent fraction of RDP proportionally increases the region over which

190

Table 13.2

TGA–DTA data for flame-retarded epoxy formulations. DTA 1st Stage

DTA 2nd Stage

DTA 3rd Stage

Sample

T10 (1C)

DT10 (1C)

Char at 800 1C (%)a

Epoxy Epoxy/RDP-95/5 Epoxy/RDP-90/10 Epoxy/RDP-85/15 Epoxy/RDP-80/20 Epoxy/RDP/I.30E95/ 0/5 Epoxy/RDP/I.30E90/5/5 Epoxy/RDP/I.30E85/10/5 Epoxy/RDP/I.30E80/15/5

364 337 336 330 319 362

0
27
28
34
45
2

0 0 0 0 3 0

150–475 150–450 150–450 150–450 150–450 150–450

390, 370; 343; 327; 322; 381,

463; (3, 4) (2) 417; (2, 2) (2) (1) 446; (3, 4)

450–650 450–650 450–690 450–690 450–690 450–650

560; 560; 545; 540; 547; 563;

338


26

3

150–450

327; 380; (2, 2)

450–650

562; (12)

–

–

336


28

2

150–450

335; (1)

450–650

548; (10)

–

–

329


35

2

150–450

323; (1)

450–650

555; (11)

–

–

Temp Range (1C)

Peak Pos.; (Peak Max.) (1C); (1Cmg
1)b

Temp Range (1C)

Peak Pos.; (Peak Max.) (1C); (1Cmg
1)b

Temp Range (1C)

Peak Pos.; (Peak Max.) (1C); (1Cmg
1)b

– – 690–750 690–750 690–750 –

– – 722; (2) 726; (3) 717; (4) –

(14) (11) (10) (11) (8) (14)

a

The silica content in the char has been accounted for. Peak Max. is the DTA response–signal value at the peak position, while Peak Pos. is the peak location with respect to temperature.

b

Chapter 13

191

Thermal and Fire Performance of Flame-Retarded Epoxy Resin

the flame retardant has a thermal stabilizing effect on epoxy resin. Additive fractions above 5% result in the enhanced char formation at temperatures above 450 1C. The overall stabilization effect (OSE) of RDP at specified percent loadings is calculated via integration of the area under the D mass percent versus temperature curve using the equation:

OSE ¼

800 X

ððmass % flame
retarded epoxy sampleT Þ

ð13:1Þ

T¼50


ðmass % control epoxyT ÞÞ

2000 0 -2000

5

10

15

% Additive fraction

Figure 13.4

OSE as a function of additive fraction in epoxy resin.

20

Epoxy/RDP/I.30E - 80/15/5

Epoxy/RDP - 80/20

Epoxy/RDP - 90/10

4000

Epoxy/RDP/I.30E - 95/0/5

6000 Epoxy/RDP - 95/5

OSE (%)

8000

Epoxy/RDP/I.30E - 85/10/5

10000

Epoxy/RDP - 85/15

Epoxy/RDP/I.30E - 90/5/5

where T is the temperature of degradation (T ¼ 50–800 1C) – the results are presented in Figure 13.4. A negative OSE value is observed for 5% RDP, which suggests that at this level the additive has an overall destabilization effect. The OSE values are positive and increase linearly with additive fraction for loadings at 10% and beyond. The anticipated thermal stabilization mechanism is as follows. At high temperatures, RDP degrades into phosphoric acid, which is subsequently converted into polyphosphoric acid. The polyphosphoric acids then promote cross-linking of polymer fragments to form a stable carbonaceous char network. Thus, catalytic char induction and condensed, or vapour phase, action of phosphorus may be implicated in the enhanced char formation at elevated temperatures.21–25 DTA curves of the unmodified epoxy and RDP flame-retarded samples are shown in Figure 13.5, and the extracted data are presented in Table 13.2. The DTA curve of the unmodified epoxy features a broad exothermic peak between 150 and 470 1C with spikes at 390 and 463 1C. This event is consistent with the

192

Temp. Difference (°C/mg)

Chapter 13 Epoxy Epoxy/RDPEpoxy/RDPEpoxy/RDPEpoxy/RDP-

12

95/ 5 90/10 85/15 80/20

8

4

0 50

Figure 13.5

150

250

350 450 Temperature (°C)

550

650

750

DTA curves for unmodified epoxy resin and RDP flame-retarded epoxy resin samples.

decomposition of epoxy chains to form a primary carbonaceous char. A pronounced exothermic peak is observed between 450 and 650 1C with a peak maximum at 560 1C and corresponds to char oxidation. RDP-containing epoxy resin formulations show similar, but less pronounced, features to those of the unmodified resin in the lower temperature regime, which suggests a possible chemical interaction. The exothermic peak value in the char oxidation stage is progressively suppressed with an increase in RDP content (Table 13.2). The depletion of the exothermicity is consistent with the enhanced char formation at elevated temperatures – less total heat is produced from oxidative thermal degradation of flame-retarded systems. Interestingly, exothermic features are observed between 690 and 750 1C for RDP-containing epoxy and correspond to the oxidation of the secondary char into a stable and final residue. The variations in DTA profiles between the epoxy resin and its RDP flame-retarded samples suggest a possible change in the degradation mechanism of the former. However, without performing further experimental work, such as evolved gas analysis (EGA), the authors cannot speculate on probable mechanistic alterations in the presence of RDP.

13.3.2.2

Effect of RDP and Nanoclay on the Thermal Stability of Epoxy Resin

The addition of RDP together with 5% I.30E to the epoxy resin is expected to change its decomposition behaviour. TG mass-loss profiles for the unmodified resin and its nanocomposites that contain 5% I.30E together with various amounts of RDP are shown in Figure 13.6(a). The addition of the nanoclay alone [epoxy–RDP–I.30E -(95/0/5)] does not improve the thermal stability of

193

Thermal and Fire Performance of Flame-Retarded Epoxy Resin 100

Epoxy Epoxy/RDP/I.30E - 95/ 0/ 5 Epoxy/RDP/I.30E - 90/ 5/ 5 Epoxy/RDP/I.30E - 85/10/ 5 Epoxy/RDP/I.30E - 80/15/ 5

(a)

Mass %

80

60

40

20

0 10

(b)

∆Mass %

5

0

-5 Epoxy/RDP/I.30E - 95/ 0/ 5 Epoxy/RDP/I.30E - 90/ 5/ 5 Epoxy/RDP/I.30E - 85/10/ 5 Epoxy/RDP/I.30E - 80/15/ 5

-10

-15 50

150

250

350

450

550

650

750

Temperature (°C)

Figure 13.6

(a) TGA curves and (b) mass-difference curves between RDP flameretarded nanocomposites and unmodified epoxy resin versus temperature in flowing air.

epoxy resin; in fact, the thermal degradation profile is shifted to lower temperatures relative to the unmodified resin and the effective OSE value is negative (Figure 13.4). However, the OSE value for epoxy–RDP–I.30E -(95/0/5) is less negative by a factor of two when compared to that of the epoxy–RDP (95/5) formulation. This shows that the nanoclay does not enhance the thermal stability of the epoxy resin, with the only effect seen at temperatures above 600 1C when most of the resin has already decomposed to leave silica. On the addition of 5% nanoclay together with 5, 10 and 15% RDP, the onset of thermal degradation, T10, is reduced in some cases by as much as 35 1C relative to that of the unmodified resin (Table 13.2). Similar observations were made for resin samples that contained RDP only (Table 13.2). While the flame-retarded nanoclay-containing

194

Chapter 13

formulations are destabilized in the lower temperature regime, at temperatures above 450 1C a stabilization effect is observed [Figure 13.6(b)]. Positive D mass percent values for temperatures above 450 1C suggest that the concomitant presence of RDP and the nanoclay promote char formation. The calculated OSE for epoxy resin containing 5% nanoclay and 5% RDP [epoxy–RDP–I.30E (90/5/5)], is 1.5 times higher than when RDP is used alone at 10% loading. Moreover, when used individually at 5% fractional loadings, both RDP and the nanoclay show an overall destabilization effect, as shown by negative OSE values (Figure 13.4). The combination of RDP and the nanoclay significantly improves the thermal stability of epoxy resin as measured by TGA, which suggests synergistic interactions – the combined stabilization effect is higher than a mere additive effect. However, increasing the weight fraction of RDP in the presence of clay destabilizes the resin in the lower temperature regime faster than the rate at which it promotes char formation, such that the OSE is effectively lowered relative to formulations that contain RDP alone at the same cumulative additive fractions, Figure 13.4. The combination of RDP and the nanoclay in epoxy resin is anticipated to be interactive, and thus promote the formation of char bonded structures. To evaluate synergistic and/or antagonistic interactions between RDP and the nanoclay, theoretical mass loss profiles were calculated from a linear combination of their individual profiles and compared to the experimental data. Figure 13.7 shows the calculated TG mass-loss profile of epoxy–RDP–I.30E (90/5/5) calculated from linear combination of a 1:1 mixture of epoxy–RDP (90/10) and epoxy–RDP–I.30E (90/0/10), experimental mass-loss profile and their mass difference as a function of temperature. These data suggest an

100 80 Mass %

60 1 40

∆ Mass %

3

Calculated

-1 20

Experimental

0

-3 50

150

250

350

450

550

650

750

Temperature (°C)

Figure 13.7

Calculated TG mass loss versus temperature for epoxy–RDP–I.30E (90/5/5), computed from a linear combination of a 1:1 mixture of epoxy– RDP (90/10) and epoxy–RDP–I.30E (90/0/10) ( . . . .), experimental mass-loss profile (—) and their mass difference (----) as a function of temperature.

195

Thermal and Fire Performance of Flame-Retarded Epoxy Resin Epoxy Epoxy/RDP/I.30E - 95/ 0/ 5 Epoxy/RDP/I.30E - 90/ 5/ 5 Epoxy/RDP/I.30E - 85/10/ 5 Epoxy/RDP/I.30E - 80/15/ 5

Temp. Difference (°C/mg)

12

8

4

0 50

150

250

350

450

550

650

750

Temperature (°C)

Figure 13.8

DTA curves for the unmodified epoxy resin and its RDP–nanoclay flame-retarded nanocomposites.

enhanced char formation between 350 and 550 1C and beyond 650 1C when RDP is used at low concentration (5%) together with the nanoclay as compared to when they are used individually. The exact mechanism of interaction between RDP and the nanoclay when interspersed in epoxy resin is not known; however, it is anticipated that these flame-retardant additives slow down or prevent depolymerization of the epoxy resin. DTA curves of unmodified epoxy and its RDP and nanoclay flame-retarded composites are shown in Figure 13.8. All samples exhibit a broad exothermic feature between 250 and 450 1C, with some samples showing spikes around 380 and 430 1C. A second exothermic feature that spans the temperature range 450– 650 1C and corresponds to the oxidation of primary char is observed for all samples. However, a notable feature is the depletion of the DTA exothermic response as a function of additive fraction, which suggests a change in the degradation mechanism that promotes char formation and thus possibly leads to the reduction in the amount of heat evolved.

13.3.3 13.3.3.1

Flammability Behaviour Limiting Oxygen Index

The LOI is a quantitative method via which the relative flame retardancy of materials can be evaluated. LOI is defined as the minimum oxygen concentration in an oxygen–nitrogen mixture that will just support flaming combustion. Table 13.3 and Figure 13.9 show LOI and DLOI (LOI of flame-retarded formulation minus LOI of unmodified epoxy resin) values of the unmodified

196

Chapter 13

Table 13.3

Vertical UL-94 and LOI data for flame-retarded epoxy formulations. Vertical UL-94 a

Sample

TBT (s)

Epoxy Epoxy–RDP (95/ 5) Epoxy–RDP (90/10) Epoxy–RDP (85/15) Epoxy–RDP (80/20) Epoxy–RDP–I.30E (95/0/5) Epoxy–RDP–I.30E (90/5/5) Epoxy–RDP–I.30E (85/10/5) Epoxy–RDP–I.30E (80/15/5)

88 7 3 13 15 110 23 8 16

Performance

LOI

HB V-0 V-0 V-1 V-1 HB V-1 V-0 V-0

27.8 30.2 30.2 28.6 29.6 30.0 27.8 27.8 27.5

a

Epoxy/RDP/I.30E - 80/15/5

Epoxy/RDP - 80/20

0.5

Epoxy/RDP - 85/15

Epoxy/RDP - 90/10

1.5

Epoxy/RDP/I.30E - 85/10/5

2.5

Epoxy/RDP/I.30E - 90/5/5

∆ LOI

3.5

Epoxy/RDP/I.30E - 95/0/5

4.5

Epoxy/RDP - 95/5

TBT, total burning time (s).

-0.5 5

10

15

20

% Additive fraction

Figure 13.9

DLOI (LOI of flame-retarded formulation minus LOI of unmodified resin) values vs. % additive fraction in epoxy resin.

epoxy resin and its flame-retarded formulations that contain RDP alone and, in some cases, together with the nanoclay I.30E. The LOI of the unmodified epoxy was determined to be 27.8. Addition of RDP at the 5% level increases the LOI value to 30.2. Increasing RDP to 10% had no effect, while beyond a 10% additive fraction resulted in reduced LOI values. These results suggest the existence of a threshold additive fraction beyond which increasing the phosphorus content in epoxy has no beneficial effect, in this case 1% phosphorus

Thermal and Fire Performance of Flame-Retarded Epoxy Resin

197

content. Phosphorus-containing formulations are expected to have improved flame resistance – it is understood that phosphoric and polyphosphoric acids promote char formation.26 When used alone in the epoxy resin, the nanoclay gives a higher LOI value [sample epoxy–RDP–I.30E (90/0/50), LOI ¼ 30.0) than the unmodified resin. However, when RDP is used in conjunction with the nanoclay, no improvement in LOI is observed. Despite the synergistic interactions predicted from TGA of epoxy–RDP–I.30E (90/5/5), from LOI analysis the combination of RDP and the nanoclay is not as effective as would have been expected (Figure 13.9). The authors envisage antagonistic interactions between RDP and the nanoclay in an attempt to improve flame resistance as measured by the LOI.

13.3.3.2

Vertical UL-94 Test

UL-94 ratings are used to describe the ease with which a polymeric material may be burned or extinguished. UL-94 vertical test results are presented in Table 13.3. The unmodified resin failed to satisfy the minimum requirements of UL-94 in the vertical testing mode, with a total burning time of 88 seconds and vigorous dripping leading to the ignition of surgical cotton. Epoxy resin formulations that contained 5 and 10% RDP achieved the V-0 rating. It is interesting that there is a remarkable increase in flame retardancy at low phosphorus content (5 and 10% RDP), while with the increase of phosphorus (15 and 20% RDP) diminishing returns are realized. A similar effect was observed from LOI measurements. Incorporation of nanoclay alone does not improve the flame retardancy of epoxy considerably, as measured by the vertical UL-94 test – the sample passed the horizontal burning test, but failed to satisfy the minimum requirements for the vertical test. Addition of RDP together with the nanoclay increases the flame resistance of the epoxy. Contrary to results observed in the LOI tests, flame resistance as measured by UL-94 increases with the amount of RDP when the nanoclay content is fixed at 5% w/w fraction. The presence of RDP together with nanoclay promotes charring at the surface during burning, and thus insulates the underlying polymeric material, which slows the mass-loss rate of decomposition products and reduces the tendency to drip. Dissimilar flame resistance results obtained from LOI and UL-94 measurements suggest that the action and efficiency of flame retardants in achieving their primary purpose is highly dependent on the heating conditions and environment.

13.3.3.3

Cone Calorimetry

13.3.3.3.1 Effect of RDP on Fire Performance of Epoxy Resin. The flammability behaviour of epoxy resin and its formulations that contain phosphorus and/or nanoclay has been assessed by cone calorimetry. The results, such as TTI, HHR, especially its peak value, and mass-loss rate are important to evaluate the fire safety of a material. The results obtained by cone calorimetry for RDP flame-retarded epoxy resin formulations are shown in Figure 13.10

198

Chapter 13 1000

Epoxy Epoxy/RDP Epoxy/RDP Epoxy/RDP Epoxy/RDP -

HRR(kW/m2)

800

95/ 5 90/10 85/15 80/20

600 400 200 0 0

50

100

150

200

250

Time (s)

Figure 13.10

HRR curves for the unmodified epoxy resin and RDP flame-retarded samples from cone calorimetry measurements at 50 kW m
2.

and the data are presented in Table 13.4. TTI, measured as the time required for the sample to burn with a sustained flame, which coincides with the onset of the HRR curves, is considerably reduced with the addition of RDP (Table 13.4). These results indicate that RDP-containing resins are thermally unstable and decompose at low temperatures to produce combustible volatiles, which are essential for a sustained flaming process, consistent with TG data shown in Figure 13.3(a). The presence of flame retardants may increase the viscosity of the unmodified resin, and so reduce the heat exchange between the samplesurface exposed to the radiant source and the bulk of the sample. This would result in a rapid increase in the surface temperature of the sample, and hence the time required for the volatiles to reach the pyrolysis temperature is significantly reduced.27 The HRR is considered to be a parameter of paramount importance in characterizing the intensity of a fire and is also related to mass-loss rate, and hence the fire growth rate. The peak HRR (PHRR) values of resin formulations that containing 5–15% RDP are significantly reduced, in some cases by 27%, when compared to that of the unmodified resin. However, the effect is less pronounced at 20% RDP additive fraction (Table 13.4). The observed reductions in PHRR suggest that the flame-retarded formulations burn to give a lowintensity flame. Products from the thermal decomposition of RDP promote cross-linking carbonization, which results in the formation of char that subsequently acts as a protective layer against heat and oxygen diffusion.28 The protective char thus prevents further decomposition of the epoxy resin to yield more residual char at the conclusion of the flaming combustion process. The fire growth rate index (FIGRA), defined as the ratio of PHRR values to the time at which they occur, tPHRR, indicates the burning propensity of a material,29 and these values are presented in Table 13.4. The control sample has

Cone calorimetry data for flame-retarded epoxy formulations at 50 kW m
2.a

Sample

TTI (s)

PHRR (kW m
2) (% red.)b

Epoxy Epoxy–RDP (95/5) Epoxy–RDP (90/10) Epoxy–RDP (85/15) Epoxy–RDP (80/20) Epoxy–RDP–I.30E (95/0/5) Epoxy–RDP–I.30E (90/5/5) Epoxy–RDP–I.30E (85/10/5) Epoxy–RDP–I.30E (80/15/5)

27 20 16 20 18 27 28 23 20

1076 995 (7) 786 (27) 787 (27) 841 (22) 915 (15) 1045 (3) 945 (12) 795 (26)

a

FIGRA (kW s
1)

THR (MJ m
2) (% red.)b

EHC (MJ kg
1)

AMLR (g s
1 m
2)

CY (%) @300 s

20 17 13 14 14 14 17 15 12

53 45(15) 42(21) 38(28) 41(23) 56(-6) 43(19) 47(11) 47(11)

21 20 20 20 19 21 20 21 22

22 20 18 18 22 21 19 16 16

16 23 30 42 31 15 21 22 27

% red., reduction in PHRR; AMLR, average mass loss rate, CY, char yield after 300 s. The coefficient of variations in cone calorimetry data are less than 10% for all parameters. b % reduction in PHRR and THR compared to the control sample.

Thermal and Fire Performance of Flame-Retarded Epoxy Resin

Table 13.4

199

200

Chapter 13

the highest FIGRA value of 20. On the addition of RDP, FIGRA values are reduced, in some cases by as much as 7 units, i.e. by 35%. This largely results from the reduction in PHRR, since tPHRR values are similar. The level of flame retardancy achieved with RDP alone as measured by the FIGRA index is notable. The reduction in FIGRA values with the addition of RDP suggests that the flame-retarded formulations burn with a lower propensity when compared to the unmodified resin. Low FIGRA values indicate delayed times to flashover, which allows enough time to evacuate and/or for fire extinguishers to arrive. THR values, as measured via cone calorimetry under a constant heat flux of 50 kW m
2 for RDP-containing epoxy resin formulations, are presented in Table 13.4. The control sample gave a THR value of 53 MJ m
2. Significant percent reductions, in the 15–28% range, in THR are observed when RDP is added to epoxy at 5–20% loadings. The addition of RDP at 15% gave the largest reduction in THR, B28%; however, addition of RDP at higher loadings, 20%, had a less pronounced effect with a reduction of 22% in THR observed. The lower THR values observed for epoxy-containing RDP are consistent with reduced fuel content from the flame-retarded samples during combustion (Table 13.4). The residual char yields significantly increase with RDP content; however, the trend is not the same as observed for char yields realized from TG experiments. This may be attributed to different heating conditions – slow heating rates are used for TG experiments, while sample temperatures rapidly evolve when samples are exposed to a high heat flux, as in cone calorimetry. The variation in the heating conditions dictates differences in flame-retardation mechanisms to yield dissimilar residual char trends from TG and cone calorimetry. The effective heat of combustion (EHC), which reflects the calorific value per unit mass of the specimen for the control sample, is 21 MJ kg
1 (Table 13.4). There is a slight reduction in EHC values with the addition of RDP. That the heat released per unit mass of volatiles is invariant with percent additive fraction for flame-retarded formulations (Table 13.4) suggests that the flame retardancy action of phosphorus does not involve flame inhibition and/or fuel dilution. The reduced amount of volatiles released from flame-retarded formulations, as shown by increased residual char yields, suggests that phosphorus is active in the condensed phase.

13.3.3.3.2 Effect of RDP and Nanoclay on Fire Performance of Epoxy Resin. HRR curves as a function of time, as measured by cone calorimetry, for RDP–nanoclay flame-retarded epoxy resin systems are shown in Figure 13.11 and the data are presented in Table 13.4. TTI does not change following addition of the nanoclay alone [epoxy–RDP–I.30E (95/0/5)]. Samples that contain RDP and the nanoclay exhibit shorter TTI values. However, their TTI values are always higher than when RDP is used individually at the same percent additive fraction. A slight reduction, 15%, in PHRR is observed for epoxy– RDP–I.30E (95/0/5), which suggests the nanoclay may somewhat slow down or

201

Thermal and Fire Performance of Flame-Retarded Epoxy Resin 1000

Epoxy Epoxy/RDP/I.30E Epoxy/RDP/I.30E Epoxy/RDP/I.30E Epoxy/RDP/I.30E -

HRR (kW/m2)

800

95/ 0/5 90/ 5/5 85/10/5 80/15/5

600

400

200

0 0

50

100

150

200

250

Time (s)

Figure 13.11

HRR curves for the unmodified epoxy resin and RDP flame-retarded nanocomposites from cone calorimetry measurements at 50 kW m
2.

prevent the production of volatiles by promoting carbonization. Interestingly, the reduction in PHRR is linearly dependent on the RDP additive fraction reaching a maximum value of 26% for epoxy–RDP–I.30E (80/15/5). Premature ignition and the observed reduction in fire intensity are indicated by lower PHRR for flame-retarded systems compared to that of the unmodified resin. This further underlines the proposition that the flame-retardation mechanism of phosphorus-based flame-retardant additives results from their decomposition at lower temperatures to yield by-products subsequently responsible for improving the fire resistance of the polymeric system at elevated temperatures. The fire safety of epoxy formulations that contain RDP and nanoclay was evaluated using FIGRA values (Table 13.4). Addition of 5–15% RDP together with 5% nanoclay reduces the FIGRA values, with a remarkable 40% reduction observed for the epoxy–RDP–I.30E (80/15/5) formulation. THR values as measured via cone calorimetry under a constant heat flux of 50 kW m
2 for RDP- and nanoclay-containing epoxy resin formulations are presented in Table 13.4. The sample that contains 5% nanoclay gave a THR value higher than that of the control sample (56 MJ m
2). The addition of clay alone [epoxy– RDP–I.30E (95/0/5)] reduced the PHRR values, but its combustion–flaming process is prolonged, which results in a higher THR value. When RDP and nanoclay are concomitantly added to the epoxy resin, significant reductions in THR values are observed. However, the percent reductions observed are not as pronounced as those obtained when RDP is used alone at the same loading fraction. This suggests that the presence of clay together with RDP may lead to antagonistic interaction, when compared to the unmodified resin, when the fire

202

Chapter 13

performance is evaluated via cone calorimetry. There is no change in the EHC for epoxy resin formulations that contain RDP and nanoclay. These results underline the proposition that phosphorus achieves its efficiency in flame retardation in the condensed phase. Contrary to the results obtained from TGA, the addition of RDP in juxtaposition with the nanoclay results in an antagonistic effect as the summed percent-gain in fire resistance is lower than when RDP and the nanoclay are used independently at the same loading fraction. However, these results are similar to what is observed in both LOI and UL-94. When used individually, RDP and the nanoclay were shown to increase LOI values, but did not improve the flame resistance when used together. When added at 5% loading fractions, RDP and the nanoclay achieved V-0 and HB classifications, respectively. On combining the two flame retardants to give a 10% cumulative additive fraction, the UL-94 classification obtained was V-1, which lies between HB and V-0. This suggests that there are no synergistic interactions between RDP and the nanoclay in reducing flammability as measured by UL-94. However, there were significant improvements in thermal stability for the resin mixture that contains 5% RDP and 5% nanoclay as measured by TGA. The incorporation of the nanoclay with RDP in epoxy resin did not give desirable results in some aspects of flame retardancy: LOI values were not improved with respect to that for the unmodified resin and also the UL-94 classification of V-0 was degraded to V-1 at low cumulative additive concentrations. The heavily cross-linked structure of thermosetting epoxy matrix may hinder the migration of nanoparticles to the pyrolysis surface, and thus restrict their effectiveness as a fire barrier.

13.4 Conclusions The interspersion of RDP in epoxy resin was shown to improve the thermal stability of the resin at elevated temperatures. The flame resistance, as measured by LOI and UL-94, of RDP-containing resin formulations is higher than for the control sample, but is invariant with the loading fraction. Formulations with RDP alone at 5 and 10% gave the best flame resistance in LOI and UL-94 experiments, while going beyond 10% RDP had an adverse effect. In cone calorimetry the addition of RDP at 10% results in remarkable reductions in both the PHRR and THR, and enhanced char formation. Addition of nanoclay alone did not improve the thermal stability of the resin, and neither did it improve its flame resistance as measured by cone calorimetry. In contrast, addition of clay together with RDP did not show an advantage. In fact the concomitant presence of RDP and the nanoclay in epoxy compromised thermal stability with the exception of the epoxy–RDP–I.30E (90/5/5) formulation. It also did not produce any significant improvement in the fire properties as measured by cone calorimetry and LOI. This may be attributed to possible antagonistic interactions between the constituent components. However, at

Thermal and Fire Performance of Flame-Retarded Epoxy Resin

203

high concentrations of RDP (10 and 15%), UL-94 ratings of V-0 are observed. The addition of a known flame retardant, RDP, and nanoclay yielded mixed results with respect to thermal stabilization and improved fire performance. Future studies will entail the use of other types of organically modified clays in an attempt to achieve efficient thermal and fire resistance for materials based on epoxy resin.

Acknowledgements We thank Betime Nuhji, Deakin University, Australia for TEM analysis. One of the authors, Charalampos Katsoulis, would like to acknowledge financial support from the Engineering and Physical Sciences Research Council (EPSRC).

References 1. G. Beyer, Nanocomposites offer new way forward for flame retardants. Plastics, Additives and Compounding, 2005, 7, 32–35. 2. S.Y. Lu and I. Hamerton, Recent developments in the chemistry of halogen-free flame retardant polymers, Prog. Polym. Sci., 2002, 27, 1661–1712. 3. G. Beyer, Nanocomposites: a new class of flame retardants for polymers, Plastics, Additives and Compounding, 2002, 4, 22–28. 4. G. Chigwada, P. Jash, D.D. Jiang and C.A. Wilkie, Synergy between nanocomposite formation and low levels of bromine on fire retardancy in polystyrenes, Polym. Degrad. Stab., 2005, 88, 382–393. 5. G. Camino and S. Lomakin, Intumescent materials, in: R.A. Horrocks, D. Price, eds Fire retardant materials, Woodhead Publishing Limited, Cambridge, UK, 2001, pp. 318–336. 6. E.T. Thostenson, C. Li and T.-W. Chou, Nanocomposites in context, Compos. Sci. Technol., 2005, 65, 491–519. 7. J. Jordan, K.I. Jacob, R. Tannenbaum, M.A. Sharaf and I. Jasiuk, Experimental trends in polymer nanocomposites – a review, Mater. Sci. Eng. A., 2005, 393, 1–11. 8. K. Putz, R. Krishnamoorti and P.F. Green, The role of interfacial interactions in the dynamic mechanical response of functionalized SWNT – PS nanocomposites, Polymer, 2007, 48, 3540–3545. 9. J. Zhang, D.D. Jiang, D. Wang and C.A. Wilkie, Styrenic polymer nanocomposites based on an oligomerically-modified clay with high inorganic content, Polym. Degrad. Stab., 2006, 91, 2665–2674. 10. G. Camino, G. Tartaglione, A. Frache, C. Manferti and G. Costa, Thermal and combustion behaviour of layered silicate–epoxy nanocomposites, Polym. Degrad. Stab., 2005, 90, 354–362. 11. P.C. Lebaron, Z. Wang and T.J. Pinnavaia, Polymer-layered silicate nanocomposites: an overview, Appl. Clay Sci., 1999, 15, 11–29.

204

Chapter 13

12. X. Kornmann, H. Lindberg and L.A. Berglund, Synthesis of epoxy-clay nanocomposites: influence of the nature of the clay on structure, Polymer, 2001, 42, 1303–1310. 13. X. Kornmann, H. Lindberg and L.A. Berglund, Synthesis of epoxy-clay nanocomposites. Influence of the nature of the curing agent on structure, Polymer, 2001, 42, 4493–4499. 14. D. Ratna, O. Becker, R. Krishnamurthy, G.P. Simon and R. Varley, Nanocomposites based on a combination of epoxy resin, hyperbranched epoxy and a layered silicate, Polymer, 2003, 44, 7449–7457. 15. A. Yasmin, J.L. Abot and I.M. Daniel, Processing of clay/epoxy nanocomposites by shear mixing, Scripta Materialia., 2003, 49, 81–86. 16. W. Liu, S.V. How and M. Pugh, Organoclay-modified high performance epoxy nanocomposites, Compos. Sci. Technol., 2005, 65, 307–316. 17. E. Kandare, G. Chigwada, D. Wang, C.A. Wilkie and J.M. Hossenlopp, Probing synergism, antagonism, and additive effects in poly(vinyl ester) (PVE) composites with fire retardants, Polym. Degrad. Stab., 2006, 91, 1209–1218. 18. B.N. Jang, M. Costache and C.A. Wilkie, The relationship between thermal degradation behavior of polymer and the fire retardancy of polymer/ clay nanocomposites, Polymer, 2005, 46, 10678–10687. 19. E.D. Weil, Additivity, synergism, and antagonism in flame retardancy, in: Flame retardancy of polymeric materials, W.C. Kuryla, A.J. Papa (ed.), Marcel Dekker, Inc, 1975, pp. 185–243. 20. E. Kandare, B.K. Kandola, J.E.J. Staggs, P. Myler. Global kinetics of thermal degradation of flame-retarded epoxy resin formulations. Polym. Degrad. Stab. 2007, 92, 1778–1787. 21. B.K. Kandola, A.R. Horrocks, P. Myler and D. Blair, Thermal characterization of thermoset matrix resins, in: G.L. Nelson, C.A. Wilkie, eds. Fire and Polymers: Materials and Solutions for Hazard Prevention. ACS Symp. Ser., Washington 2001, 797, p. 344–360. 22. B.K. Kandola, R.A. Horrocks, P. Myler and D. Blair, New developments in flame retardancy of glass-reinforced epoxy composites, J. Appl. Poly. Sci., 2003, 88, 2511–2521. 23. J. Green, Phosphorus containing flame retardants, In: Fire retardancy of polymeric materials, A.F. Grand, C.A. Wilkie, ed. Marcel Dekker, Inc. 2000, pp. 147–170. 24. F. Samyn, S. Bourbigot, S. Duquesne and R. Delobel, Effect of zinc borate on the thermal degradation of ammonium polyphosphate, Thermochimica Acta, 2007, 456, 134–144. 25. M. Jimenez, S. Duquesne and S. Bourbigot, Characterization of the performance of an intumescent fire protective coating, Surf. Coatings Tech., 2006, 201, 979–987. 26. S. Gaan and G. Sun, Effect of phosphorus and nitrogen on flame retardant cellulose: A study of phosphorus compounds, J. Anal. Appl. Pyrolysis, 2007, 78, 371–377. 27. S. Nazare´, T.R. Hull and B. Biswas, Study of relationship between rheological and flammability properties of flame retarded poly(butylene

Thermal and Fire Performance of Flame-Retarded Epoxy Resin

205

terephthalate) containing nanoclays. Materials in 18th Annual BCC conference on Flame Retardancy of Polymeric Materials, 21–23 May, 2007, Stamford, USA. 28. J. Deng and W. Shi, Synthesis and effect of hyperbranched (3-hydroxyphenyl) phosphate as a curing agent on the thermal and combustion behaviours of novolac epoxy resin, Eur. Polym. J., 2004, 40, 1137–1143. 29. S. Bourbigot, E. Devaux and X. Flambard, Flammability of polyamide6/clay hybrid nanocomposite textiles, Polym. Degrad. Stab., 2002, 75, 397–402.

Intumescents

CHAPTER 14

Porosity Estimates of Intumescent Chars by Image Analysis J.E.J. STAGGS Energy Resources Research Institute, University of Leeds, Leeds, LS2 9JT, UK

14.1 Introduction Intumescent chars are attractive candidates for incorporation into fire protection systems in many different scenarios, including steel protection and more general fire retardancy applications. The production of a robust, coherent, highly porous char is desirable from a thermal insulation viewpoint. Although such structures have low density by definition, they also have low thermal conductivity. This is not the complete picture, however. Consider an insulating layer of material of thickness l, thermal conductivity k, subject to an external heat flux q_ 00 on one exposed surface. The temperature drop across the layer is of the order DTBq_ 00 l=k. Thus, we see that for a given heating situation, the temperature on the unexposed face, and hence the performance of the insulator, is governed by the conduction heat transfer coefficient k0 ¼ k=l. Ideally, k should be small enough (or l large enough) to ensure that the temperature of the unexposed face remains acceptably low. This observation of the importance of k 0 is of pivotal significance to intumescent chars: their very nature is to have high volumetric expansion (large l ) and hence high porosity (which implies low k).

Fire Retardancy of Polymers: New Strategies and Mechanisms Edited by T Richard Hull and Baljinder K Kandola r Royal Society of Chemistry 2009 Published by the Royal Society of Chemistry, www.rsc.org

209

210

Chapter 14

Correctly assigning values for the density and specific heat capacity of a porous solid is unproblematic: if the subscript g denotes a pore property and the subscript s denotes a solid matrix property, then the composite density and specific heat capacity are given, respectively, by r ¼ jrg þ ð1
jÞrs ; c ¼ lcg þ ð1
lÞcs Ecs

ð14:1Þ

where j is the porosity (the ratio of the pore volume to the total volume) and l is the mass fraction of pores, which will be small compared with one for most chars of interest. In general, thermal conductivity is a decreasing function of porosity, but the exact dependence depends on, among other things, the shape and distribution of the pores within the char in relation to the direction of heat transfer. This point is illustrated by the direct numerical simulations shown in Figure 14.1. Here we see the results of a finite-element calculation of the effective thermal conductivity of a porous solid with two different orientations of plate-like pores. The overall porosity for both cases is fixed at 25%, with ks/kg ¼ 100, where k is thermal conductivity. In the first case the pores are aligned such that their long axes are perpendicular to the direction of heat transfer and in the second case the pores are aligned with their long axes parallel to the direction of heat transfer. The shaded contours represent temperature isotherms. Note that the case where the pores’ long axes are normal to the direction of heat transfer has the best insulating effect. The reason for this is that orientation of the low thermal conductivity pores is such that they more effectively disrupt the conductive pathway in the direction of heat transfer.

(a) k / ks = 0.31

Figure 14.1

(b) k / ks = 0.7

Direct numerical simulation to illustrate the effect of pore orientation on effective thermal conductivity, j ¼ 25%, ks/kg ¼ 100.

211

Porosity Estimates of Intumescent Chars by Image Analysis

In general it is possible to place bounds on the dependence of effective thermal conductivity on porosity. For any porous solid consisting of a solid matrix of thermal conductivity ks, with gas-filled voids of conductivity kg, it is easy to demonstrate that the effective thermal conductivity k must be between two extremes obtained by considering two thermal resistors, either in series or in parallel:1
 1 k 1 ks  1
j 1
;k¼ kj þ ð1
jÞ ks k kg

ð14:2Þ

For some cases it is possible to find explicit expressions for the dependence of k on j. For spherical voids of fixed radius and low porosity, the modified Maxwell expression1,2 k 2jð1
kÞ þ 1 þ 2k E ks 2k þ 1
jð1
kÞ

ð14:3Þ

gives reasonable results (see Figure 14.2 for an illustration). Another expression due to Bruggeman2 assumes that the voids consist of spheres with an infinite range of radii. In this case k/ks is given implicitly by the equation k=ks
1=k ðk=ks Þ1=3 ð1
1=kÞ

¼1
j

ð14:4Þ

1 0.9

(parallel)

0.8 0.7 Maxwell's approximation

k / ks

0.6 0.5

Effect of pore size & shape for fixed porosity

0.4 0.3 0.2 (series)

0.1 0 0

0.2

0.4

0.6 Porosity

Figure 14.2

Effective thermal conductivity bounds (k ¼ 10).

0.8

1

212

Chapter 14

When k is large and j is not close to one, it may be shown that the solution of Bruggeman’s expression gives
1=3
2 k 1
ð1
jÞ3=2 1 þO Eð1
jÞ1=2 þ ks k 2kð1
jÞ

ð14:5Þ

Here O(1/k)2 means terms that involve powers of 1/k of order two and higher. Hence a major step towards estimating the thermal conductivity of a char may be made if the following key facts are known:  The porosity of the char.  The shape and/or orientation of the pores with reference to the direction of heat transfer.  The distribution of pore sizes.

14.2 Pore-Finding Algorithm Consider a two-dimensional (2D) rectangular digital image of a char section that has been segmented into pure black and white. The char pores appear as collections of black pixels and the solid matrix of the char as white pixels. Let the segmented image contain a total of n black pixels. Furthermore, assume that the pixels are labelled in some consistent way and that we have a list P ¼ fPi gni¼1 of the labels of all n black pixels. In relation to the segmented image, a pore is defined as a collection of connected black pixels. Connections can be defined in one of two main ways, depending on whether a connection between two pixels is defined as a common edge or a common vertex (Figure 14.3). Common vertex connection results in fewer, larger pores than does common edge connection. Once we have decided on the connection definition, the algorithm to identify a pore is straightforward. The goal is to assemble a list of pixels p ¼ fpj gm j¼1  P that are connected and hence form a pore. This may be done efficiently as follows:  Starting with an empty pore list p, define a seed for the pore that corresponds to the first pixel label in P and add it to the pore list p.  Delete the seed from P. Common Edge Connection: Pixels 1 and 2 share a side and so are connected, but not pixel 3, so there are two pores

1 2 3

Figure 14.3

Common Vertex Connection: Pixels 1 and 2 share common vertices and pixels 2 and 3 share a common vertex, so all are connected and there is one pore.

Pixel connection definitions.

Porosity Estimates of Intumescent Chars by Image Analysis

Test Case 1: Porosity = 0.185

Figure 14.4

213

Test Case 2: Porosity = 0.26

Test cases for pore-finding algorithm.

 Search through P and add to p any members that are connected to any members of p.  Each time a connected member of P is found, delete it from P. At the end of the search, P will be depleted by the pixels comprising the pore list p, which gives a list of pixel labels that form a connected pore. This search is then repeated to find the next pore and so on until P is empty. To check the pore-finding algorithm, two negative images were prepared to show white pores of known porosity in a black matrix (Figure 14.4). In both cases the algorithm correctly identified the pores and the overall porosity. For each pore, an equivalent pore radius (EPR) and aspect ratio (AR) is defined. These correspond, respectively, to the radius of a circle with the same area as that of the pore and the ratio of the maximum height to maximum width of a pore (Figure 14.5).

14.3 Relationship Between Area Porosity and Volume Porosity We define the area porosity, j2D , as the ratio of pore area to total area as determined in a 2D section through a char, and the volume porosity, j3D , as the true porosity of the 3D char, i.e. the ratio of pore volume to total volume. Given a porous char, it is not immediately obvious that the area porosity, obtained from a section of the char, will be the same as the volume porosity. To see this, consider a single sphere circumscribed by a cube. Taking a section parallel to one of the faces of the cube, we see that the intersected sphere will produce a circle and so the resulting area porosity will be anything in the range

214

Chapter 14

EPR: circles of the same area as the pore

B

Aspect Ratio =

B A

A

Figure 14.5 Definition of EPR and AR.

Section Plane

r 2 − z2

{

r

z

3D sphere

Figure 14.6

2D Section

Difference between area and volume porosity.

from 0 to p/4, depending where the section is taken, whereas the volume porosity is always p/6 (Figure 14.6). With this in mind, consider the relationship between area and volume porosity for a porous char with a large number of pores whose centres are randomly distributed throughout the char. A thin section of thickness dz and side l through the char intersects a certain number n of pores and produces a 2D plane of intersected pore slices (Figure 14.7).  2 , where A is the average area of the The area porosity is given by j2D ¼ nA=l intersected pores. Now, since the total volume of pores intersected by the 2  it follows that the volume porosity is j3D ¼ nAdz=ðl  section is nAdz, dzÞ ¼ j2D . Hence, if the pore centres are uniformly distributed, we would expect the area

Porosity Estimates of Intumescent Chars by Image Analysis

215

3D pore

Section Plane

Figure 14.7

2D Section

Intersected pore distribution.

and volume porosities to be identical for a representative section taken through the char. To illustrate this result, Monte Carlo simulations were carried out. In the simulations, spheres were randomly allocated positions within a cube so that their (x, y, z) co-ordinates were uniformly distributed and their radii followed a given distribution. A section through the middle of the cube was taken and the area porosity computed. This was then compared with the volume porosity, and the results are shown in Figure 14.8. This figure was computed using uniformly distributed radii, although the same result was obtained for other distributions.

14.4 Relationship Between 2D and 3D Pore Distributions It is important to realise that a probability density of radii obtained from a 2D section will not be the same as the true 3D probability density. However, the key question is whether or not we can reconstruct the 3D distribution from the 2D data. To answer this, we must first investigate the relationship between the 2D and 3D distributions. Consider again a porous char with pores that consist of randomly distributed spheres within a cube of side l. Now, if the z-ordinates are uniformly distributed and the probability density function (PDF) of the radii is f ð^rÞ, where ^r ¼ r=l,

216

Chapter 14 0.3 r ~ U(0,0.025)

y=x

Volume Porosity

0.25 0.2 0.15 0.1 0.05 0 0

Figure 14.8

0.05

0.1

0.15 Area Porosity

0.2

0.25

0.3

Results of Monte Carlo simulations for randomly distributed spheres.

since r and z are independent, it may be shown3 that the PDF of radii in the 2D section f ð2DÞ ðrÞ is given in terms of the true PDF as

f

ð2DÞ

ðrÞ ¼ 2r

Z1

r

f ð^ rÞd^r ð^ r2


r2 Þ1=2

ð14:6Þ

By way of illustration, we now consider two test cases. In each test case Monte Carlo simulations were carried out whereby spheres were randomly allocated positions within a cube so that their (x, y, z) co-ordinates were uniformly distributed. The PDF of intersected sphere radii was then computed and the calculation repeated 50 times in order that an average PDF could be found and compared to an appropriate expression obtained from Equation (14.6).

14.4.1

Test Case 1 (Identical Spheres)

If the 3D spheres all have the same radius, r0 say, then it can be shown from above that the distribution of radii from a 2D section will have a PDF given by f ð2DÞ ðrÞ ¼

r ð1
r2 Þ1=2

ð14:7Þ

where r ¼ r=r0 . The graph in Figure 14.9 shows the results of the Monte Carlo simulation compared with the predicted PDF.

217

Porosity Estimates of Intumescent Chars by Image Analysis 0.35 Monte Carlo Results Prediction

0.3

Relative Frequency

0.25 0.2 0.15 0.1 0.05 0 0

0.2

0.4

0.6

0.8

1




Figure 14.9

14.4.2

Comparison between expected frequency distribution and Monte Carlo results for test case 1.

Test Case 2 (Spheres with Uniformly Distributed Radii)

If the 3D spheres have uniformly distributed radii, ^rBUð0; 1Þ, then it can be shown from above that the distribution of radii from a 2D section will have a PDF given by   pffiffiffiffiffiffiffiffiffiffiffiffiffi 1 1 þ 1
r2 f ð2DÞ ðrÞ ¼ 2r ln r

ð14:8Þ

where r ¼ r/l. The graph in Figure 14.10 shows the results of the Monte Carlo simulation, carried out as in the first test case above, and compared with the predicted PDF.

14.5 Construction of 3D Distributions from 2D Distributions We have seen above that the distribution of radii from a 2D section is not the same as the distribution of radii of the 3D spheres from which the section was taken. In particular, we have the result that for a random distribution of spheres of identical radius, the distribution of radii in a 2D section is given by the probability density, Equation (14.7). The important question here is whether or not we can construct the PDF for the 3D distribution from the PDF of

218

Chapter 14 0.08 0.07

Relative Frequency

0.06 0.05 0.04 0.03 0.02

Monte Carlo Results Prediction

0.01 0 0

0.2

0.4

0.6

0.8

1




Figure 14.10

Comparison between expected frequency distribution and Monte Carlo results for test case 2.

the 2D section. In fact, as we shall see presently, it is possible to use the relationship shown in Section 14.4 to construct the 3D distribution of pore radii from the 2D distribution. In principle, this involves the solution of a difficult Volterra integral equation of the second kind.4 In practice it is rarely possible to achieve the exact solution (especially using experimental data for the 2D distribution) and so a numerical method must be used. Suppose that the 2D distribution is given in terms of (n+1) discrete points ð2DÞ fi ¼ f ð2DÞ ðri Þ, i ¼ 0,1, . . . , n. Then we seek to construct f on the same set of points using a simple numerical rule for the integral. The only difficulty lies in evaluation of the integral in the neighbourhood of r^ ¼ r. Now, let e be small. Since 9 8rþ Z1 > ; i ¼ n
2; > > 1=2 2 > < ðriþ1
r2i Þ " # Si ¼ n
1 > X > f f iþ1 j >r h > > : i ðr2
r2 Þ1=2 þ 2 2 2 1=2 j¼iþ2 ðrj
ri Þ i iþ1

ð14:12Þ

ð14:13Þ

and Si ¼ 0, i > n
2. To test the method, we attempt to reconstruct the 3D PDF from a 2D section obtained from a known distribution of spheres. We may then compare the reconstructed PDF with the known distribution. So, if a 2D section is generated from randomly distributed spheres with uniformly varying radius, i.e. r B U(0,l) where l is the maximum radius in the 3D distribution, then the expected 3D PDF should simply be f(r) ¼ 1, 0 r r r 1, where r ¼ r/l. Therefore, using the 2D PDF   pffiffiffiffiffiffiffiffiffiffiffiffiffi 1 1 þ 1
r2 f ð2DÞ ðrÞ ¼ 2r ln r

ð14:14Þ

the numerical method above was applied to construct the expected 3D PDF. The results are shown in Figure 14.11 and indicate clearly that the expected U(0,1) distribution for r is recovered.

14.6 Analysis of a Real Char Section The diagram in Figure 14.12 shows a segmented image of a section through a real intumescent char (supplied by International Paint Ltd), with an expansion ratio of approximately 10, along with a representative region selected for analysis. Note that in this case the char formed numerous very large pores close to the top surface – which appears in the segmented image as a dark irregular band close to the top surface. The z-direction is taken out of the plane of the figure, the x-direction is horizontal and the y-direction is vertical. The char expanded in such a way as the structure in the x- and z-directions is broadly similar.

220

Chapter 14 1.4

2D Distribution

1.2 1 PDF

Constructed 3D Distribution 0.8 0.6 0.4 0.2 0 0

0.2

0.4

0.6

0.8

1




Figure 14.11

Reconstructed 3D distribution using test data for the 2D distribution.

Figure 14.12

Sample segmented image of a char section.

221

Porosity Estimates of Intumescent Chars by Image Analysis 0.0025

Relative Frequency

0.0020

0.0015

0.0010

0.0005

0.0000 1500

2000

2500

3000

3500

4000

4500

x-ordinate of pore

Figure 14.13

Frequency distribution of x-ordinates of pores.

Individual pores were determined from the pixel data using the algorithm of Section 14.2 and the frequencies of the x-ordinates of their centroids are shown in Figure 14.13. These data indicate strongly that the x-ordinates are uniformly distributed. Given the expansion of the char, it is reasonable to assume that the char structure is similar in the x- and z-directions. It may well not be in the y-direction, because this corresponds to the main direction of heat transfer as the char was formed. Hence, it seems reasonable to deduce from the x-ordinate data that the z-locations of the pores are uniformly distributed and so the conditions of the result established in Section 14.3 apply. Hence, we may use the result of Section 14.4 to attempt to construct the 3D distribution of EPR from the 2D data. The graph in Figure 14.14 shows the relative frequencies of pore radii (triangles) from the 2D segmented image. The squares on the plot correspond to the application of the method of Section 4.5 to obtain the reconstructed 3D distribution. Note that the reconstructed 3D PDF has more noise than the 2D distribution – a direct consequence of the method – but it is still useable. Note also that a small number of very small pores were found from the image analysis, but these have been deleted as they are likely to be below the resolution of the scanner used to produce the original segmented image. A log–log plot of the 3D frequency distribution strongly suggests a PDF of the form

f ðrÞ ¼


 n
1 r
n ; r  rmin rmin rmin

ð14:15Þ

222

Chapter 14 r / max(r) 0

0.2

0.4

0.6

0.8

1

1

Relative Frequency

0.1 0.01 2D Data

0.001 0.0001 3D reconstruction 0.00001 0.000001

Figure 14.14

Comparison between empirical 2D distribution (triangles) and reconstructed 3D distribution (squares).

ln(r/r min) 0

0.5

1

1.5

2

2.5

3

0

ln(Relative Frequency)

-2 -4 -6 -8 -10 y = -3.6103x - 1.2024 -12

R2 = 0.934

-14

Figure 14.15

Confirmation of exponential distribution.

and is shown in Figure 14.15, where the exponent n has been found as approximately 3.6. The fitted frequency obtained from Equation (14.15) is shown by the solid smooth curve in Figure 14.14. Finally, the graph in Figure 14.16 shows the frequency distribution of the pore AR. This figure indicates that most pores are oblate – being compressed along the axis in the direction of heat transfer. This is actually a desirable property, as demonstrated in the introduction, since for the same porosity a char that consists of oblate, disc-like pores (where the axis of the disc is parallel to the direction of heat transfer) will have a lower thermal conductivity than a

223

Porosity Estimates of Intumescent Chars by Image Analysis 0.008

Relative Frequency

0.007 0.006 0.005 0.004 0.003 0.002 0.001 0 0

0.2

0.4

0.6

0.8

Aspect Ratio

Figure 14.16

Frequency distribution of pore AR.

char that consists of perfectly spherical pores.5,6 The Bruggeman model5,6 for oblate spheroids, flattened in the direction of heat transfer, of negligible thermal conductivity (kg E 0), may be appropriate for this type of char. Here, the composite thermal conductivity is given by an expression of the form k ¼ ð1
jÞ1=ð1
F Þ ks

ð14:16Þ

where F is a shape factor of the spheroid (F ¼ 1/3 for a sphere, 1/2 for a cylinder). Note that Equation (14.5) agrees with this for F ¼ 1/3 as k - N.

14.7 Conclusion The analysis in this chapter is concerned with interpreting details about the 3D distribution of pores within a char using data obtained from an image of a 2D section through the char. The primary reason for doing this is to obtain relevant details to enable an appropriate correlation to be used to model the thermal conductivity of the char. Also, the analysis provides a tool such that the frequency distributions of EPR may be estimated and so new correlations may be developed that assume realistic pore distributions. It transpires that if the distribution of the z-ordinates of the pores is uniform (where z is measured normal to the plane of the 2D section), then the following results are valid: 1. The volume porosity is the same as the area porosity. 2. The 2D distribution of equivalent pore radii frequencies is not the same as the 3D distribution. In fact, there is a relationship between the 2D EPR

224

Chapter 14 (2D)

distribution and the 3D equivalent pore distribution. If f (r) is the PDF of the 2D distribution, where r ¼ r/rmax (r is the EPR and rmax the maximum EPR), and f(r) is the PDF of the 3D distribution, then:

f

ð2DÞ

ðrÞ ¼ 2r

Z1

r

f ð^ rÞd^r ð^r2
r2 Þ1=2

ð14:17Þ

3. The relationship in result (2) above may be used to obtain the 3D EPR distribution from the corresponding 2D distribution. Analysis of an actual section through a char reveals that the x-ordinates of the pore centroids are uniformly distributed. By symmetry, we would broadly expect a section in the x–y plane to have a similar structure to that of a section in the y–z plane. From this, it follows that if the x-ordinates are uniformly distributed, we would expect that the z-ordinates of the centroids of the 3D pores would also be uniformly distributed. Hence, the condition for the three results above should be met in practice. Naturally, this could be confirmed by an analysis of a section taken in the y–z plane. Analysis of a 2D section of a real intumescent char leads to the conclusion that the PDF of the 3D EPR is of exponential form.

Acknowledgements The author is grateful to Drs Paul Jackson and Rachel Butler of International Paints Ltd for permission to use the segmented image of a real char section.

References 1. J.E.J. Staggs, Modelling the thermal conductivity of porous materials using thermal resistor networks, Fire Safety Journal, 2002, 37, 107. 2. I.I. Kantorovich and E. Bar-Ziv, Heat transfer within highly porous chars: a review, Fuel, 1999, 78, 279. 3. P.G. Hoel, Introduction to Mathematical Statistics. Wiley, 5th Edition, y10 (1984). 4. A.J. Jerri, Introduction to Integral Equations with Applications. Wiley, 2nd Edition, y3 (1999). 5. F. Cernuschi, S. Ahmaniemi, P. Vuoristo and T. Ma¨ntlya¨, Modelling of thermal conductivity of porous materials: application to thick barrier coatings, J. Eur. Ceram. Soc., 2004, 24, 2657. 6. B. Schultz, Thermal conductivity of porous and highly porous materials, High Temp. High Press., 1981, 13, 649.

CHAPTER 15

Efficient Modelling of Temperatures in Steel Plates Protected by Intumescent Coating in Fire J.F. YUAN AND Y.C. WANG School of Mechanical, Aerospace, and Civil Engineering, The University of Manchester, M60 1QD, UK

15.1 Introduction Intumescent coating is designed to expand and form a thick, porous charred layer when exposed to heat in fire. The charred layer insulates the underlying substrate by providing a physical barrier. The applications of intumescent coating are wide-ranging and the demands for this material have significantly increased in recent years, especially in the civil engineering area to protect building structures from fire attack. However, the mechanisms that determine the fire-resistant properties of intumescent coating are not well understood yet, because of their highly complex physical and chemical natures. This hampers the application of intumescent coating in performance-based fire engineering. At present, building fire resistance design is largely based on the ‘‘standard fire condition’’, and the thermal properties (mainly thermal conductivity) of an intumescent coating are obtained under the standard temperature–time relationship.1 However, in a natural fire condition in buildings, fire development varies depending upon such factors as the amount of

Fire Retardancy of Polymers: New Strategies and Mechanisms Edited by T Richard Hull and Baljinder K Kandola r Royal Society of Chemistry 2009 Published by the Royal Society of Chemistry, www.rsc.org

225

226

Chapter 15

combustible materials, ventilation and construction lining materials. For conventional fire protection materials, their thermal properties can be pre-determined because they are only temperature dependent. Unlike the conventional fire protection materials, intumescent coating will behave differently according to the applied fire condition, coating thickness and protected structures. Intumescent coating is a reactive material, so the temperature–thermal properties relationship of an intumescent coating may be entirely different under different realistic fire conditions, and any set of temperature–thermal properties relationship may only be applicable to a very narrow range of fire conditions. A key to widen the applicability of any predictive method is to ensure that it captures the essential features of intumescent coating that govern coating behaviour in fire. So far a number of researchers have developed predictive models with different degrees of complexity. At the simplest level, intumescent coating is treated in the same way as the more conventional non-reactive fire protection materials, and different effective thermal conductivity values of intumescent coating are used for different conditions. This is clearly not satisfactory because it will not be possible to extrapolate the results to different applications. At the most complex level, a few researchers2 have attempted to model the intumescence process at the microscopic level, with a detailed consideration of hydrodynamics, heat transfer and chemical reactions related to individual bubble nucleation, production, movement and burst. This method of simulation is computationally prohibitive, as it requires a large amount of input data. So far it has had limited success, being used only to demonstrate the intumescence process during the very early stage of heating. An alternative third method would be to consider only the essential features of the global chemical decomposition and physical behaviour of intumescent coating. This has been followed by a number of researchers using different assumptions.3,4 The focus of this study is on the temperatures of the protected steel, rather than the detailed behaviour of intumescent coating itself. Therefore, the primary objective was to develop a practical method to extract the basic thermal properties of intumescent coating that are independent of fire or heat. Therefore, we adopted the third approach and the framework of Di Blasi and Branca,4 but with a few revised assumptions that are pointed out in Section 15.2. According to Di Blasi and Branca, intumescent coating decomposition is modelled in three global steps and their reaction rates are governed by the Arrhenius equation. The intumescence process is simply represented by expansion of the intumescent thickness. As is shown in Section 15.3, even though the model adopted in this study may be considered simple, it still requires a large number of input data for the intumescent coating. It may be difficult to obtain all of these input data precisely. Therefore, an important objective was to identify the input parameters that will have significant influence on the calculated steel temperatures, by conducting a sensitivity study to investigate the influences of different input parameters on steel temperatures.

Efficient Modelling of Temperatures in Steel Plates

227

Figure 15.1 Illustration of energy conservation within an infinitesimal layer of intumescent coating.

15.2 Mathematical Modelling For a steel plate protected by intumescent coating exposed to fire, it is acceptable to model the assembly as a one-dimensional entity. By discretizing the intumescent coating into a number of layers, such as that shown in Figure 15.1, finite difference method (FDM) can be used to solve the corresponding equations. From energy conservation, the change in heat flow conducted through a discretized layer of intumescent coating is equal to the change in internal energy of the coating, which gives (see the section Nomenclature for the symbol usage):   @ @T @T l ¼ ðms Cs þ mg Cg Þ @x @x @t @ðxrg Þ @ms þ Cs T þ Cg T @t  @t _ @ðgTÞ @mr Cg þ Dh þ @x @t

ð15:1Þ

The left-hand side of Equation (15.1) is the conducted heat to the intumescent coating; the first term on the right-hand side (RHS) is the heat increase caused by changes in coating temperature. The second and third terms on the RHS are heat increases resulting from changes in the gas and solid masses, respectively. The fourth term on the RHS describes convective heat loss due to gas movement and the last term on the RHS is heat release from the coating. Equation (15.1) theoretically applies to the whole intumescent char; however, it is practically solved within discretized layers, as aforementioned, because of the non-linearity of coating properties throughout the thickness. To enable

228

Chapter 15

quantification of Equation (15.1), it is necessary to derive the various masschange terms, which are determined by the decomposition process. The Arrhenius law is used to describe decomposition process, giving:   Ej Kj ¼ Aj exp
; j ¼ 1; 2; 3 ð15:2Þ